Adaptive solid-state luminescent phosphors

ABSTRACT

The absorbance or emission wavelength of composite materials comprising a transition metal doped shell disposed over a rare earth doped core and a functionalizable group on the surface of the transition metal doped shell can change upon subjection to a carboxylic acid. This method of changing the absorbance or emission wavelength of a composite material can be used to identify counterfeit currency using an ink comprising a composite material.

CROSS-REFERENCE TO RELATED APPLICATIONS

The present application claims priority to U.S. Provisional ApplicationNo. 62/562,594, filed Sep. 25, 2017, which is incorporated by referenceherein in its entirety.

BACKGROUND OF THE INVENTION

As of 2005, the US Department of the Treasury estimated that roughly 60%of all US currency was held abroad, roughly $450 billion (Cameron, etal., The Use and Counterfeiting of United States Currency Abroad, Part3; Department of the Treasury: 2006). As such, most of this is held incountries with volatile political and economic conditions due to thestability of the dollar over local currencies. Of the currency abroad,it is estimate that 1 in 10,000 bills are counterfeit. The currentaccepted method of detection is with a commercially available detectionpen (Carmeli, Method of Detecting Counterfeit Paper Currency. 1991);however, this can easily be defeated using common household items, suchas hairspray.

Since the mid-90s, holograms, among other methods, have played animportant role in anti-counterfeit technologies (Hardwick, et al., Adv.Mater. 2001, 13, 980-984). However, these holograms are imbedded intothe document and can be lifted from the surface to be reused elsewhere.More recently, anti-counterfeit technologies have focused ondownconversion dyes and nanoparticles which emit visible light whenexcited with <400 nm excitation, concealing images and patterns understandard conditions (Liu, et al., Nanoscale 2011, 3, 4804-4810). Theseinks can be easily purchased online for the emission of blue, green, redand IR light with a commercially available UV light (Ldp Llc—Maxmax.Comwww.maxmax.com).

Alternatively, there are reports of upconversion inks using rare earth(RE) doped nanostructures which convert IR light to visible (Liu, etal., Nanoscale 2011, 3, 4804-4810; You, et al., Nanoscale 2015, 7,4423-4431). Upconversion phosphors require a detailed knowledge of thesystem in order to accurately replicate, as the luminescent spectrum ishighly dependent on the type and concentration of RE dopants used (Mi,et al., Sci. Rep. 2016, 6, 22545). One of the most common upconversioncombinations is Er³⁺/Yb³⁺ (Dorman, et al., J. Phys. Chem. C 2012, 116,10333-10340). In this pairing the Er³⁺ is the emitter ion and Yb³⁺ actsas a sensitizer, increasing the amount of light absorb and transferringthe energy to the active ions. As the Yb³⁺ concentration is increased,the red-to-green emission ratio increases and results in a “redder”appearance (Mi, et al., Sci. Rep. 2016, 6, 22545). Additionally, it ispossible to produce green and blue upconversion luminescence bysubstituting the Er³⁺ with other RE ions, such as Tm³⁺ (blue) and Ho³⁺(green). (Yi, et al., J. Mater. Chem. 2005, 15, 4460-4464).

During the incorporation of RE elements into a crystal, the f-orbitalsare shielded from external interactions and split into discrete energies(Sun, et al., Annual review of physical chemistry 2015, 66, 619-642).Upon splitting, the parity selection rules begin to break down and toallow for previously forbidden intra-orbital transitions (Judd, PhysicalReview 1962, 127, 750; Ofelt, The Journal of Chemical Physics 1962, 37,511-520). This phenomenon allows for distinct transitions which resultin a systematic, luminescent “fingerprint” unique to each RE. Thisfingerprint is nearly identical for all crystal hosts due to the stable3+ oxidation state (Dieke, et al., Spectra and Energy Levels of RareEarth Ions in Crystals; Interscience Publishers New York, 1968; Vol. 5).As such, RE luminescence has been the key component in transformativetechnologies over the past half century, including lasers, opticaldisplays, fiber optic communications, and biological imaging (Zhou, etal., Chemical Reviews 2015, 115, 395-465; Zhou, et al., Nat Nano 2015,10, 924-936). However, it is also this 3+ oxidation state, andsimilarity in chemical properties, that make the separation of thesematerials difficult, requiring environmentally hazardous chemicals. Thisfact has led to a material dependence on China for the manufacturing ofour current technologies (Bauer, et al., Critical Materials Strategy.Energy, D. o., Ed. Washington D.C., 2010; p 166).

Over the past decade there has been a push for the development ofalternative luminescent materials to prevent this dependence. Onetypically avoided class of materials is the first-row transition metal(TM) elements. While this group of elements is known to have an intenserange of colors, they are also highly susceptible to the coordinationand strength of the bonded ligand (Orgel, J. Chem. Soc. 1952, 4756-4761;Kiang, et al., Quantum Electronics, IEEE Journal of 1965, 1, 295-298).Crystal field theory was developed to describe the d-orbitalhybridization and splitting based on metal oxidation state, ligand type,and arrangement of bonds (Van Vleck, Physical Review 1932, 41, 208-215;Griffith, and Orgel, Quarterly Reviews, Chemical Society 1957, 11,381-393). These parameters can be described with a single term derivedby Giulio Racah to define field strengths (Racah's B variable) (Racah,Physical Review 1942, 61, 186; Racah, Physical Review 1942, 62, 438;Racah, Physical Review 1943, 63, 367; Racah, Physical Review 1949, 76,1352). From crystal field theory, it is possible to predict the opticalproperties of TM ions in solids and as molecular complexes resultingfrom the d-orbital splitting, Δ. While the splitting and associatedenergy level filling is dictated by the number and strength of ligands,the most common structures can be divide into a triply degenerate(t_(2g)-d_(xy), d_(xz), d_(yz)) and a double degenerate (e_(g)-d_(x) ₂_(-y) ₂ , d_(z) ₂ ) energy states. To account for the available opticaltransitions, Tanabe and Sugano created a set energy diagrams for theexcited states of a six-fold coordinated (octahedral) molecular complexwhich are normalized to the crystal field strength (Δ) (Tanabe andSugano, Journal of the Physical Society of Japan 1954, 9, 753-766;Tanabe and Kamimura, Journal of the Physical Society of Japan 1958, 13,394-411; Tanabe and Sugano, Journal of the Physical Society of Japan1954, 9, 766-7791 Tanabe and Sugano, Journal of the Physical Society ofJapan 1956, 11, 864-877). These plots can be used to extract crystalfield strengths from optical measurements by correlating measuredtransitions to expected transitions. However, Tanabe-Sugano diagramshave limitations in predicting optical properties of solids due tochemical stabilities and electrostatic interactions required for crystalgrowth. In order to engineer the electrostatic interactions within acrystal, complex stoichiometries are required.

The challenge with engineering electrostatic interactions lies in thenature and size of the crystal. Specifically, crystal field theoryrelies on the optical properties of bulk crystals (>100s of nm) (Orgel,J. Chem. Soc. 1952, 4756-4761; Weakliem, The Journal of Chemical Physics1962, 36, 2117-2140). However, as the push for smaller, more powerfulelectronics continues, active material size must decrease. This oftenresults in unwanted energy transfer effects as observed in luminescentmaterials (Dorman, et al., The Journal of Physical Chemistry C 2012,116, 10333-10340; Tanaka, et al., Journal of Luminescence 2000, 87,472-474; Bhargava, et al., Physical Review Letters 1994, 72, 416).Typically, energy transfer is ∝ R⁻⁶, where R is the distance between theexcited state and the sink, as defined by Forster and Dexter in the1960s (Forster, Naturwissenschaften 1946, 33, 166-175; Dexter, TheJournal of Chemical Physics 1953, 21, 836-850). Based on thisrelationship, the energy transfer rapidly decreases within ˜10 nm, asseen in experimental results (Dorman, et al., The Journal of PhysicalChemistry C 2012, 116, 10333-10340).

In more conventional solid-state physics, extremely high pressures areused to break crystal symmetry and distort electron densities (Li, etal., Inorganic Chemistry 2016, 55, 6770-6775). Similarly, large electricfields can be applied to manipulate electron density. However,application of these external fields is difficult without costlyfabrication techniques (Weste and Eshraghian, Principles of Cmos VlsiDesign; Addison-Wesley New York, 1985; Vol. 188). Recently, the abilityto control charge injection across a TiO₂-organic molecule interface wasdemonstrated by applying a dipole to the surface (Goh, et al., Journalof Applied Physics 2007, 101, 114503-114503). The surface dipolereorganized the electron density at the interface, acting as a weakdiode, and regulated charge transport into the metal oxide. Thisapproach offers an elegant method to reversibly tune interfacialelectron density, with dipole field strengths proportional to R⁻⁵(Moreno, et al., International journal of quantum chemistry 1994, 52,829-835). By carefully designing the nanostructures, TM dopants have thepotential to play a major role in solid-state luminescence instead oftheir current application as secondary RE sensitizer (Dan, et al.,Materials Letters 2015, 150, 76-80; Tian, G., et al., Advanced Materials2012, 24, 1226-1231).

There remains a need in the art for nanophosphors that emit specific andtunable wavelengths. This invention addresses this unmet need.

SUMMARY OF THE INVENTION

In one aspect, the present invention relates to a composite materialcomprising a rare earth doped core; a transition metal doped shelldisposed over the core; and a functionalizable group on the surface ofthe transition metal doped shell. In one embodiment, the compositematerial forms a shape selected from the group consisting of acore-shell nanoparticle, a nanowire, a nanorod, and a thin film disposedover a substrate. In one embodiment, the composite material is acore-shell nanoparticle. In one embodiment, the rare earth doped corecomprises β-NaYF₄. In one embodiment, the rare earth doped corecomprises at least one rare earth selected from the group consisting ofEr, Yb, Tb, Tm, and Ho. In one embodiment, wherein the transition metaldoped shell comprises TiO₂. In one embodiment, the transition metaldoped shell comprises at least one transition metal selected from thegroup consisting of V, Cr, Mn, Fe, Co, Ni, Cu, and Bi. In oneembodiment, the transition metal doped shell comprises Ni. In oneembodiment, the functionalizable group is a hydroxide group. In oneembodiment, the transition metal doped shell comprises Bi. In oneembodiment, at least one of the rare earth doped core and the transitionmetal doped shell comprises YVO₄.

In another aspect, the present invention relates to an ink comprisingthe inventive composite material. In aspect embodiment, the presentinvention relates to a QR code comprising said ink.

In another aspect, the present invention relates to a method of changingthe absorbance or emission spectrum of a nanoparticle, the methodcomprising: providing a nanoparticle having a rare earth doped core, atransition metal doped shell, and at least one surface functionalizablegroup; and treating the nanoparticle with at least one carboxylic acid;wherein the absorbance or emission spectrum of the nanoparticle ischanged upon treatment with the carboxylic acid. In one embodiment, thestep of treating the core-shell nanoparticle with at least onecarboxylic acid comprises: treating the nanoparticle with a firstcarboxylic acid; and treating the nanoparticle with a second carboxylicacid. In one embodiment, the carboxylic acid is selected from the groupconsisting of para-(fluorosulfonyl)benzoic acid, para-nitrobenzoic acid,para-cyanobenzoic acid, para-bromobenzoic acid, benzoic acid,para-methoxybenzoic acid, and para-aminobenzoic acid.

In another aspect, the present invention relates to a method ofidentifying counterfeit currency, the method comprising the steps of:providing an ink having a core-shell nanoparticle with surfacefunctionalizable groups; applying the ink during the minting ofauthentic currency; treating a currency sample with a solutioncomprising a carboxylic acid; and exposing the currency sample to UVlight; wherein the treatment with carboxylic acid changes the emissionwavelength of the core-shell nanoparticle. In one embodiment, thecore-shell nanoparticle comprises a transition metal doped shelldisposed over a rare earth doped core. In one embodiment, the transitionmetal doped shell comprises at least one transition metal selected fromthe group consisting of V, Cr, Mn, Fe, Co, Ni, Cu, and Bi.

BRIEF DESCRIPTION OF THE DRAWINGS

The following detailed description of preferred embodiments of theinvention will be better understood when read in conjunction with theappended drawings. For the purpose of illustrating the invention, thereare shown in the drawings embodiments which are presently preferred. Itshould be understood, however, that the invention is not limited to theprecise arrangements and instrumentalities of the embodiments shown inthe drawings.

FIG. 1, comprising FIGS. 1A to 1C, depicts the nanoparticle design anddual luminescence effects. FIG. 1A depicts the core-shell nanophosphorconstruction. FIG. 1B is an inset depicting the functionalization of thesurface of the nanophosphor. FIG. 1C depicts a $100 bill under UVexcitation before (top) and after (bottom) application of a chemicalagent.

FIG. 2, comprising FIGS. 2A and 2B, depicts the effect of crystalstructure and chemical treatment on absorption. FIG. 2A plots how theCr³⁺ absorption changes in a Ruby and an Emerald due to crystalstructure. Insets show the respective geometries of the transition metalcomplexes. FIG. 2B depicts possible Er³⁺ upconversion mechanisms afterchemical treatment for red and green emission.

FIG. 3, comprising FIGS. 3A to 3D, depicts yttrium core-shell andcore-multishell nanostructure characterization. FIG. 3A is a plotshowing Y—Y and Y—O bond lengths as a function of annealing temperaturein the conversion of Y(OH)₄ to Y₂O₃. FIG. 3B is a TEM image showing theprecise Y₂O₃|Yb₂O₃ interface of a core-shell nanoparticle. FIG. 3Cdepicts the CIE color coordinates of a core-multi shell phosphor withcontrolled energy transfer between layers. FIG. 3D is an inset showingthe emission spectrum and the rare earths responsible for each peak.

FIG. 4, comprising FIGS. 4A to 4D, depicts the soft XAS (CAMD), x-rayphotoemission spectroscopy (XPS), and UV-Vis absorption of a thin filmof Ni²⁺-doped TiO₂ with and without surface functionalization. FIG. 4Ais a plot of the XAS of the O K edge of an undoped thin film, a dopedthin film, and doped thin films functionalized with para-substitutedbenzoic acids. Inset shows the NiO-like (decreased cation oxidationstate) and TiO₂-like (increased cation oxidation state) regions. FIG. 4Bis a plot of the XAS of the Ni LIII/LII edges of a doped thin film anddoped thin films functionalized with para-substituted benzoic acids.FIG. 4C is a plot of the XPS scans of an undoped thin film, a doped thinfilm, and doped thin films functionalized with para-substituted benzoicacids, indicating a change in electron density between the four thinfilms. FIG. 4D depicts the shift in Ni²⁺ absorption withfunctionalization in comparison to the NaYF₄:Er³⁺|Y(OH)₃ system.

FIG. 5 depicts the simulated and experimental spectrum of an anataseTiO₂ crystal O K edge. Insert shows the 7-atom TiO₂ cluster.

FIG. 6, comprising FIGS. 6A and 6B, depicts the TEM and EELS ofFe₃O₄|MnFe₂O₄ nanoparticles. FIG. 6A is a TEM of the nanoparticles. FIG.6B is an EELS of the nanoparticles, showing the location of Fe (green)and Mn (purple) ions.

FIG. 7, comprising FIGS. 7A and 7B, depicts a time dependent densityfunctional theory (TD-DFT) model for the study of Ni²⁺-doped TiO₂. FIG.7A depicts the 123 atom model for a “bulk” solid material. FIG. 7Bdepicts the model functionalized with a para-substituted benzoic acid.

FIG. 8 depicts a comparison between the predicted Ni²⁺ and Er³⁺ energylevels. The dashed boxes highlight overlap between Er³⁺ emission levelsand possible Ni²⁺ transitions for energy transfer.

FIG. 9, comprising FIGS. 9A and 9B, demonstrates the weak field tuningof the optoelectronic properties of transition metal doped solid hosts.FIG. 9A shows the t_(2g):e_(g) intensity ratio in Ni L_(II) edge,observed to be a function of the dipole moment of the ligand. Thisspectral difference with the dipoles indicates the change in thehybridization of the Ni—O bond in TiO₂:Ni. FIG. 9B shows that theinfluence of the ligand on the core-levels of Ni (2p) is not observed inXPS due to the strong atomic multiplet coupling.

FIG. 10, comprising FIGS. 10A to 10D, depicts the effect of weak fieldtuning on the host. FIG. 10A shows the broadening of the peaks in the O(2p)-Ti/Ni (4sp) hybridization region in the O K edge. FIG. 10B showsthe dampening of the e_(g) shoulder peak in the Ti L_(III) edge. Thesetwo plots indicate non-cubic structural distortion upon Ni doping inanatase TiO₂ crystal. This geometric distortion is attributed to theoxygen vacancies that are formed in the lattice due to chargecompensation. FIG. 10C shows The energy difference between the t_(2g)and e_(g) hybridization peaks (ΔE) in the lower energy region of O Kedge. FIG. 10D shows the full width at half-maximum (FWHM) of the e_(2g)peak in the Ti L_(II) edge, which is observed to increase with theadsorbate dipole moment, suggesting that the difference in the crystalfield splitting energy (10 Dq) of the TiO₂:Ni films is due to the changein the hybridization of the metal-oxygen bond.

FIG. 11, comprising FIGS. 11A and 11B, depicts a comparison between themeasured and modeled Ni L_(III/II) edge in TiO₂:Ni films functionalizedwith para-benzoic acid (BZA) ligands (R=NO₂, 3.8 D and R=NH₂, −4.5 D).FIG. 11A depicts the experimental results. FIG. 11B depicts the resultsfrom computational modeling experiments.

FIG. 12, comprising FIGS. 12A and 12B, show a preliminary model of a Tianatase cluster. FIG. 12A shows the Ti₉O₃₈H₆₀ anatase cluster preparedby covalently embedded procedure, wherein the boundary O atoms arepassivated by pseudo H atoms to replicate the bulk Ti atoms. FIG. 12Bshows the simulated XANES spectra of Ti L edge using Real time (RT)-TimeDependent Density Functional Theory (TDDFT) approach. The doublet of thee_(g) peak in the Ti L_(III) edge is attributed to the lowering of theTi symmetry from O_(h) to D_(2d) in anatase crystal structure (insert).

FIG. 13, comprising FIGS. 13A to 13C, demonstrates the ability to tuneabsorbance by manipulation of dopant and host materials. FIG. 13A showsthe UV-Vis absorption spectra of NaYF₄:Er³⁺|Y(OH)₃:Ni²⁺ core-shellnanoparticles (NPs) and demonstrates that the Ni²⁺ absorption spectracan be shifted by ˜70 nm with weak chemical dipoles. FIG. 13B shows theNi²⁺ sensitized Er³⁺ emission in NaYF₄:Er|Y(OH)₃:Ni core-shell NPs,demonstrating higher luminescence intensities for Er³⁺ emission at 575nm. FIG. 13C shows the power dependence of luminescence intensity ofgreen and blue emissions in Ni²⁺ sensitized Er³⁺ follows similar trendas that of Yb sensitized Er emission, demonstrating that Ni²⁺ is asuitable replacement for Yb³⁺ sensitizer in upconversion phosphors. Allof these optical characterization results point to the ability of Ni²⁺to sensitize Er³⁺ emission for tunable optical responses via adaptiveabsorption of the Ni²⁺ ion.

FIG. 14, comprising FIGS. 14A and 14B, shows the effect of aging time onsegregation. TiO₂:Ni (15 mol %) nanoparticles are quantified using XRDand UV-Vis to demonstrate the structure-processing-propertyrelationship. FIG. 14A shows the XRD pattern of TiO₂:Ni (15 mol %) NPsthat are annealed directly after drying or aged after drying show aclear difference with respect to the formation of NiO. FIG. 14B showsthat the absorption peaks are broadened or shifted in Ni doped TiO₂ whencompared to the NiO segregated TiO₂ nanoparticles. These results showthat the optoelectronic properties of the highly doped TiO2-basedmaterials can be tuned by varying the processing parameters.

FIG. 15, comprising FIGS. 15A and 15B, shows the effect of aging time onsegregation due to the hydroxyl cluster formation. The chemical changesduring the aging of TiO₂:Ni (15 mol %) dried powders is observed fromFTIR and time-resolved UV-Vis absorption measurements. FIG. 15A showsthe FTIR spectra of TiO₂:Ni (15 mol %) dried powders that are aged inair for 48 h, which indicate an increase in the surface hydroxylconcentration (3500 cm⁻¹) accompanied by dampening of the C—H alkanestretches around 2900 cm⁻¹. FIG. 15B shows that the UV-Vis absorptionpeaks are systematically shifted to lower wavelengths upon aging in air,indicating an increase in the ligand field strength (10 Dq) due to thebonding of the cation to more electronegative groups (—OH). Both ofthese results suggest that amorphous metal hydroxide clusters are formedupon aging in air, which further transform into metal oxide clustersafter annealing.

FIG. 16, comprising FIGS. 16A and 16B, show a similar phenomenon ofdopant incorporation and dopant segregation in TiO₂:Co (15 mol %)nanoparticles. FIG. 16A is an XRD pattern of TiO₂:Co²⁺ (15 mol %) NPsand shows that rapid annealing can lock the dopants in the host latticewithout forming a segregated Co₃O₄ phase. FIG. 16B depicts the UV-Visabsorption spectra of Co doped TiO₂ and Co₃O₄ segregated TiO₂nanoparticles. These results further reinforce thestructure-processing-property relationship in highly doped TiO₂materials, wherein the crystal structure is controlled by varying theannealing rate to obtain modified optical properties.

FIG. 17, comprising FIGS. 17A and 17B, demonstrate the minimalsegregation propensity in Fe and Cr dopants. FIG. 17A shows the XRDpattern of TiO₂:Fe²⁺ (15 mol %) NPs. FIG. 17B shows the XRD pattern of)TiO₂:Cr³⁺ (15 mol %) NPs.

FIG. 18, comprising FIGS. 18A and 18B, show the luminescence propertiesof YVO₄:Eu³⁺YVO₄:Bi³ core-shell nanoparticles. FIG. 18A shows the PLexcitation spectra of YVO₄:Eu³⁺ (5 mol %), YVO₄:Eu³⁺ (5 mol %), Bi³⁺ (6mol %), and YVO₄:Eu³⁺ (5 mol %)|YVO₄:Bi³⁺ (6 mol %) core-shell NPs. FIG.18B shows the shift in the excitation spectra of YVO₄:Eu³⁺ (5 mol%)|YVO₄:Bi³⁺ (6 mol %) core-shell NPs ˜30 nm using weak chemical dipoles(para-substituted benzoic acid ligands). This property of tunableexcitation can be exploited to selectively tune the emission intensitiesfor select applications.

FIG. 19, comprising FIGS. 19A and 19B, compares the phosphorescentlifetimes of core-shell NPs and surface-modified core-shell NPs. FIG.19A shows the phosphorescent lifetime of YVO₄:Eu³⁺ (5 mol %), YVO₄:Eu³⁺(5 mol %), Bi³⁺ (6 mol %), and YVO₄:Eu³⁺ (5 mol %)|YVO₄:Bi³⁺ (6 mol %)core-shell NPs. FIG. 19B shows the phosphorescent lifetimes ofpara-substituted YVO₄:Eu³⁺ (5 mol %)|YVO₄:Bi³⁺ (6 mol %) core-shell NPs.There is a slight reduction in the lifetimes of the ligand bondedcore-shell NPs.

FIG. 20, comprising FIGS. 20A and 20B, depicts AFM images of annealedthin films. FIG. 20A depicts the AFM image of a TiO₂ thin film. FIG. 20Bdepicts the AFM image of a TiO₂:Ni (15 mol %) thin film. AFM scans wereperformed on Agilent 4500 microscope in contact mode. The scanning areawas 2×2 μm² and the images were processed using Gwyddion software forcalculating roughness and other parameters. The roughness of pure TiO₂thin film was obtained as 2-3 nm whereas the Ni doped film was about40-50 nm.

FIG. 21, comprising FIGS. 21A and 21B, depict imaging and quantificationof TiO₂:Ni (15 mol %) thin films. FIG. 21A shows the UV-Vis absorptionspectra, inset shows the O_(h) coordination of Ni²⁺ and electronictransitions according to theory. FIG. 21B shows the GI-XRD pattern ofTiO₂ thin film indexed to anatase TiO₂ (JCPDS#12-1272).

FIG. 22 is a Tanabe-Sugano diagram of Ni²⁺ (3d⁸) in O_(h) symmetry. Theboxed electronic transitions are spin allowed transitions as the spinmultiplicity remains same both in the ground and excited states. Slopesof all the spin allowed transitions are calculated as shown in theadjacent Tanabe-Sugano diagram. With the observed experimental values ofE, the crystal field parameters Dq=890.4 cm⁻¹ and B=780.27 cm⁻¹ areobtained by solving the ³A₂-³T₂ (³F) and ³A₂-³T₁ (³P) two simultaneousequations. To check the accuracy of the calculated parameters(Dq/B=1.14), the λ of other two transitions ³A₂-¹E and ³A₂-³T₁ (³F) aretheoretically computed. The curvy pattern in the absorption spectra from600-800 nm is attributed to the deviation from perfect octahedralenvironment, which results in the splitting of the ³T₁ (³F) excitedstate.

FIG. 23 is a plot showing the GI-XRD pattern of TiO₂:Ni (15 mol %) filmdeposited on an Si(100) substrate and annealed at 450° C. for 2 h. Thecrystal structure was observed to be destroyed probably because of theformation of new NiTiO₃ phase and non-uniformity of the film.

FIG. 24, comprising FIGS. 24A to 24D, shows Ni-dense regions of aTiO₂:Ni film. FIG. 24A is a bright field HRTEM image of the TiO₂:Nifilms showing TiO₂ and Ni dense regions (highlighted in yellow). FIG.24B is an EELS chemical map of Ni. FIG. 24C is an EELS chemical map ofTi. FIG. 24D is an EELS chemical map of O. The EELS chemical mapsdemonstrate the high and low concentration Ni phases. The scale bars inFIG. 24 correspond to 5 nm.

FIG. 25, comprising FIGS. 25A and 25B, depicts SAED and FFT patterns ofTiO₂:Ni (15 mol %) film on Si substrate. The lattice spacings wereextracted as 3.54 Å and 1.91 Å for (101) and (020) planes of TiO₂ (red),and 2.12 Å for (002) plane of NiTiO₃ (green). FIG. 25A depicts the SAEDpattern. FIG. 25B depicts the FFT pattern.

FIG. 26, comprising FIGS. 26A and 26B, depicts TEM imaging of TiO₂:Ni.FIG. 26A is a dark field TEM image. FIG. 26B is an HRTEM image ofTiO₂:Ni (15 mol %) superimposed with TiO₂ (red) and NiTiO₃ (green)planes.

FIG. 27, comprising FIGS. 27A to 27D, depicts EELS imaging and spectraof three regions of Ni: I (Ni poor), II (Ni present), and III (Ni rich)in TiO₂:Ni (15 mol %) films. FIG. 27A is an EELS image indicating thethree regions. FIG. 27B is an EELS spectra of Ti L_(III,II) in the threehighlighted regions. FIG. 27C is an EELS spectra of O K in the threehighlighted regions. FIG. 27D is an EELS spectra of Ni L_(III,II) in thethree highlighted regions. The inability to identify Ni L edge in Nipoor region is attributed to the high detection limit of thespectrometer

FIG. 28, comprising FIGS. 28A to 28D, depicts X-ray photoelectronspectra for pure TiO₂ and TiO₂:Ni (15 mol %) thin films. FIG. 28Adepicts the results from survey scans. FIG. 28B depicts the XPS of Ti2p. FIG. 28C depicts the XPS of O 1s. FIG. 28D depicts the XPS of Ni 2p.A shift in intensities from the O₁ (lattice oxygen) to the O₂(sub-lattice oxygen) with Ni doping is observed and attributed to oxygendefects for charge compensation.

FIG. 29, comprising FIGS. 29A to 29D, depicts the Ti 2p and O 1sdeconvoluted spectra for pure TiO₂ and TiO₂:Ni film. FIG. 29A depictsthe deconvoluted Ti 2p for pure TiO₂. FIG. 29B depicts the deconvolutedTi 2p for TiO₂:Ni (15 mol %) films. FIG. 29C depicts the deconvoluted O1s for pure TiO₂. FIG. 29D depicts the deconvoluted O 1s for TiO₂:Ni (15mol %) films.

FIG. 30 depicts the FTIR spectra for the surface modified TiO₂ films.

FIG. 31 depicts the XPS scans of Ni 2p in TiO₂:Ni (15 mol %) film aftersurface modification with benzoic acid (BZA) ligand (offset).

FIG. 32, comprising FIG. 32A to 32C, depicts the deconvoluted Ni 2p XPSspectra of various thin films. FIG. 32A is the deconvoluted Ni 2p XPSspectrum of TiO₂:Ni (15 mol %). FIG. 32B is the deconvoluted Ni 2p XPSspectrum of TiO₂:Ni|NO₂-BZA. FIG. 32C is the deconvoluted Ni 2p XPSspectrum of TiO₂:Ni|NH₂-BZA.

FIG. 33, comprising FIGS. 33A to 33D, depicts deconvoluted XANESspectra. FIG. 33A is a deconvoluted XANES spectrum of the O K edge inligand bonded TiO₂:Ni films. FIG. 33B is a deconvoluted XANES spectrumof the Ti L edge in a TiO₂:Ni film. FIG. 33C is a deconvoluted XANESspectrum of the Ti L edge in a TiO₂:Ni|NO₂-BZA film. FIG. 33D is adeconvoluted XANES spectrum of the Ti L edge in a TiO₂:Ni|NH₂-BZA film.

FIG. 34 depicts a Ni K Edge XANES of ligand bonded TiO₂:Ni films. Ni Kedge was measured in the HEXAS beam line in fluorescence mode with a Ge(111) crystal as monochromator. Be crystal was employed as the detectorat a distance of 8 cm from the sample. The data was calibrated to K edgeof NiO powder.

FIG. 35, comprising FIGS. 35A to 35E, depicts experimental and modeledNi LIII/II edge spectra of surface-modified TiO₂:Ni (15 mol %). FIG. 35Adepicts the experimental Ni LIII/II edge spectra of surface modifiedTiO₂:Ni (15 mol %) films showing a change in the branching ratio(t_(2g)/e_(g)) in L_(II) edge as a function of the ligand. FIG. 35Bshows a comparison between calculated and experimental data for the NiL_(III/II) edge spectra using computer program CTM4XAS for O_(h)/NH₂-BZA(top) and D_(4h)/NO₂-BZA (bottom) Ni²⁺ symmetry. FIG. 35C showsball-and-stick representations of octahedral (O_(h)) symmetry, squareplanar symmetry, and distorted octahedral symmetry and are shown tohighlight the change in local symmetry. The calculated spectra wereshifted by a constant (1 eV) to match the absolute values of theexperimental spectra.

FIG. 36, comprising FIGS. 36A to 36C, depicts the interplay between thesurface dipole, electronic states, p-d hybridization, and the crystalfield splitting energy of Ni²⁺ 3d orbitals. FIG. 36A shows theelectronic density of states in TiO₂:N. FIG. 36B shows the energy of theNi 3d states. FIG. 36C shows the crystal field splitting energy (10 Dq)of the solid with the ligand.

FIG. 37, comprising FIGS. 37A and 37B, depicts TGA and DSC plots ofamorphous TiO₂:Ni (15 mol %) powders showing differences in weight lossand heat flow with respect to the aging time of the dried powders. FIG.37A is a TGA plot of an amorphous TiO₂:Ni (15 mol %) powder. FIG. 37B isa DSC plot of an amorphous TiO₂:Ni (15 mol %) powder.

FIG. 38 is a schematic showing the hydration of amorphous TiO₂:Ni matrixupon aging in air. Upon annealing, the crystalline phases of NiO andTiO₂ co-exist due to the stabilization of Ti(OH)₄ clusters in thehydrated matrix.

FIG. 39 is an HRTEM of the rapidly annealed TiO₂:Ni (15 mol %) NPs,demonstrating homogeneous doping of Ni in TiO₂ without forming anysegregated phases.

DETAILED DESCRIPTION

The invention relates to adaptive nanophosphors that demonstratespecific wavelengths that are tunable through transition metal dopingand surface treatments.

Definitions

It is to be understood that the figures and descriptions of the presentinvention have been simplified to illustrate elements that are relevantfor a clear understanding of the present invention, while eliminating,for the purpose of clarity, many other elements found in the art relatedto thin film production, nanoparticles, doped materials, and the like.Those of ordinary skill in the art may recognize that other elementsand/or steps are desirable and/or required in implementing the presentinvention. However, because such elements and steps are well known inthe art, and because they do not facilitate a better understanding ofthe present invention, a discussion of such elements and steps is notprovided herein. The disclosure herein is directed to all suchvariations and modifications to such elements and methods known to thoseskilled in the art.

Unless defined otherwise, all technical and scientific terms used hereinhave the same meaning as commonly understood by one of ordinary skill inthe art to which this invention belongs. Although any methods, materialsand components similar or equivalent to those described herein can beused in the practice or testing of the present invention, the preferredmethods and materials are described.

As used herein, each of the following terms has the meaning associatedwith it in this section.

The articles “a” and “an” are used herein to refer to one or to morethan one (i.e., to at least one) of the grammatical object of thearticle. By way of example, “an element” means one element or more thanone element.

“About” as used herein when referring to a measurable value such as anamount, a temporal duration, and the like, is meant to encompassvariations of ±20%, ±10%, ±5%, ±1%, or ±0.1% from the specified value,as such variations are appropriate.

Throughout this disclosure, various aspects of the invention can bepresented in a range format. It should be understood that thedescription in range format is merely for convenience and brevity andshould not be construed as an inflexible limitation on the scope of theinvention. Accordingly, the description of a range should be consideredto have specifically disclosed all the possible subranges as well asindividual numerical values within that range. For example, descriptionof a range such as from 1 to 6 should be considered to have specificallydisclosed subranges such as from 1 to 3, from 1 to 4, from 1 to 5, from2 to 4, from 2 to 6, from 3 to 6 etc., as well as individual numberswithin that range, for example, 1, 2, 2.7, 3, 4, 5, 5.3, 6 and any wholeand partial increments therebetween. This applies regardless of thebreadth of the range.

DESCRIPTION

In one aspect, the invention relates to a composite material comprisinga transition metal doped shell disposed over a rare earth doped core,wherein the thin film possesses surface functionalizable groups.

The composite material can be of any shape known to those of skill inthe art. In one embodiment, the composite material is a nanowire,wherein the transition metal doped shell forms a sheath over the rareearth doped core. In another embodiment, the composite material is amultilayered thin film disposed over a substrate. In another embodiment,the composite material is a core-shell nanoparticle with a rare earthdoped core and a transition metal doped shell.

In one embodiment, the rare earth doped core comprises a transparentmaterial which effectively separates rare earth ions. In one embodiment,the rare earth doped core comprises NaYF₄, NaGdF₄, LiYF₄, Gd₂O₃, LaGaO₃,Sc₂O₃, Y₃Al₅O₁₂ (YAG), YVO₄, Y(OH)₃, YF₃, CaF₂, HfO₂, ZrO₂, TiO₂, orLu₂O₃. In one embodiment, the rare earth doped core comprises NaYF₄. Inone embodiment, the rare earth doped core comprises β-NaYF₄.

In some embodiments, the rare earth doped core comprises at least onerare earth element selected from the group consisting of cerium (Ce),dysprosium (Dy), erbium (Er), europium (Eu), gadolinium (Gd), holmium(Ho), lanthanum (La), lutetium (Lu), neodymium (Nd), praseodymium (Pr),promethium (Pm), samarium (Sm), scandium (Sc), strontium (Sr), terbium(Tb), thulium (Tm), ytterbium (Yb) and yttrium (Y). In one embodiment,the rare earth doped core comprises Er. In one embodiment, the rareearth doped core comprises Yb. In one embodiment, the rare earth dopedcore comprises Ho. In one embodiment, the rare earth doped corecomprises Tm. In one embodiment, the rare earth doped core comprises Tb.In one embodiment, the rare earth doped core comprises more than onerare earth element.

In one embodiment, the rare earth element content of the rare earthdoped core is between 0 mol % and 100 mol %. In one embodiment, the rareearth element content of the rare earth doped core is between 0 mol %and 75 mol %. In one embodiment, the rare earth element content of therare earth doped core is between 0 mol % and 50 mol %. In oneembodiment, the rare earth element content of the rare earth doped coreis between 0 mol % and 25 mol %. In one embodiment, the rare earthelement content of the rare earth doped core is between 0.1 mol % and 25mol %. In one embodiment, the rare earth element content of the rareearth doped core is between 0.2 mol % and 25 mol %. In one embodiment,the rare earth element content of the rare earth doped core is between0.3 mol % and 25 mol %. In one embodiment, the rare earth elementcontent of the rare earth doped core is between 0.4 mol % and 25 mol %.In one embodiment, the rare earth element content of the rare earthdoped core is between 0.5 mol % and 25 mol %. In one embodiment, therare earth element content of the rare earth doped core is between 0.5mol % and 10 mol %. In one embodiment, the rare earth element content ofthe rare earth doped core is about 0.5 mol %. In one embodiment, therare earth element content of the rare earth doped core is about 1.0 mol%. In one embodiment, the rare earth element content of the rare earthdoped core is about 2.0 mol %. In one embodiment, the rare earth elementcontent of the rare earth doped core is about 3.0 mol %. In oneembodiment, the rare earth element content of the rare earth doped coreis about 4.0 mol %. In one embodiment, the rare earth element content ofthe rare earth doped core is about 5.0 mol %. In one embodiment, therare earth element content of the rare earth doped core is about 6.0 mol%. In one embodiment, the rare earth element content of the rare earthdoped core is about 10 mol %. In one embodiment, the rare earth elementcontent of the rare earth doped core is about 15 mol %. In oneembodiment, the rare earth element content of the rare earth doped coreis about 20 mol %.

In some embodiments, the rare earth doped core comprises anything thatemits light. For example, in some embodiments, the rare earth doped corecomprises a first row transition metal such as titanium (Ti), vanadium(V), chromium (Cr), manganese (Mn), iron (Fe), cobalt (Co), nickel (Ni),copper (Cu), and zinc (Zn). In one embodiment, the rare earth doped coredoes not comprise a rare earth element. In one embodiment, the rareearth doped core comprises more than one material that emits light. Inone embodiment, the rare earth doped core comprises a rare earth elementand a transition metal.

In one embodiment, the rare earth doped core is less than 100 nm indiameter. In one embodiment, the rare earth doped core is less than 80nm in diameter. In one embodiment, the rare earth doped core is lessthan 60 nm in diameter. In one embodiment, the rare earth doped core isless than 50 nm in diameter. In one embodiment, the rare earth dopedcore is less than 40 nm in diameter. In one embodiment, the rare earthdoped core is less than 30 nm in diameter. In one embodiment, the rareearth doped core less than about 20 nm in diameter. In one embodiment,the rare earth doped core is less than about 10 nm in diameter.

The rare earth doped core can be fabricated using any method known tothose of skill in the art, including, but not limited to,hydrothermal/solvothermal methods, molten salt methods, sol-gel,co-precipitation, colloidal distribution, and thermal decompositionmethods.

In one embodiment, the transition metal doped shell comprises atransparent material known to those in the art. In one embodiment, thetransition metal doped shell comprises a metal or metalloid. In oneembodiment, the transition metal doped shell comprises an oxide,sulfide, selenide, or fluoride. Exemplary transparent materials include,but are not limited to, TiO₂, Al₂O₃, Be₃Al₂SiO₆, Y₂O₃, Y(OH)₄, Y(OH)₃,YVO₄, Yb₂O₃, NiO, SiO₂, CdSe, La₂O₃, Lu₂O₃, ZnO, Sc₂O₃, ZrO₂, and HfO₂.In one embodiment, the transition metal doped shell comprises TiO₂. Inone embodiment, the transition metal doped shell comprises Y₂O₃. In oneembodiment, the transition metal doped shell comprises SiO₂.

In one embodiment, the transition metal doped shell comprises at leastone transition metal. Exemplary transition metals include, but are notlimited to, titanium (Ti), vanadium (V), chromium (Cr), manganese (Mn),iron (Fe), cobalt (Co), nickel (Ni), copper (Cu), zinc (Zn), zirconium(Zr), niobium (Nb), molybdenum (Mo), technetium (Tc), ruthenium (Ru),rhodium (Rh), palladium (Pd), silver (Ag), cadmium (Cd), hafnium (Hf),tantalum (Ta), tungsten (W), rhenium (Re), osmium (Os), iridium (Ir),platinum (Pt), gold (Au), mercury (Hg), and bismuth (Bi). In oneembodiment, the transition metal doped shell comprises a transitionmetal selected from the group consisting of Ti, V, Cr, Mn, Fe, Co, Ni,Cu, and Zn. In one embodiment, the transition metal doped shellcomprises Ni. In one embodiment, the transition metal doped shellcomprises Cr. In one embodiment, the transition metal doped shellcomprises Cu. In one embodiment, the transition metal doped shellcomprises Fe. In one embodiment, the transition metal doped shellcomprises Mn. In one embodiment, the transition metal doped shellcomprises Co. In one embodiment, the transition metal doped shellcomprises V. In one embodiment, the transition metal doped shell doesnot include a rare earth. In one embodiment, the transition metal dopedshell comprises Bi.

In one embodiment, the transition metal element content of thetransition metal doped shell is between 0 mol % and 100 mol %. In oneembodiment, the transition metal element content of the transition metaldoped shell is between 0 mol % and 75 mol %. In one embodiment, thetransition metal element content of the transition metal doped shell isbetween 0 mol % and 50 mol %. In one embodiment, the transition metalelement content of the transition metal doped shell is between 0 mol %and 25 mol %. In one embodiment, the transition metal element content ofthe transition metal doped shell is between 0.1 mol % and 25 mol %. Inone embodiment, the transition metal element content of the transitionmetal doped shell is between 0.2 mol % and 25 mol %. In one embodiment,the transition metal element content of the transition metal doped shellis between 0.3 mol % and 25 mol %. In one embodiment, the transitionmetal element content of the transition metal doped shell is between 0.4mol % and 25 mol %. In one embodiment, the transition metal elementcontent of the transition metal doped shell is between 0.5 mol % and 25mol %. In one embodiment, the transition metal element content of thetransition metal doped shell is about 0.5 mol %. In one embodiment, thetransition metal element content of the transition metal doped shell isabout 1.0 mol %. In one embodiment, the transition metal element contentof the transition metal doped shell is about 2.0 mol %. In oneembodiment, the transition metal element content of the transition metaldoped shell is about 3.0 mol %. In one embodiment, the transition metalelement content of the transition metal doped shell is about 4.0 mol %.In one embodiment, the transition metal element content of thetransition metal doped shell is about 5.0 mol %. In one embodiment, thetransition metal element content of the transition metal doped shell isabout 6.0 mol %. In one embodiment, the transition metal element contentof the transition metal doped shell is about 10 mol %. In oneembodiment, the transition metal element content of the transition metaldoped shell is about 15 mol %. In one embodiment, the transition metalelement content of the transition metal doped shell is about 20 mol %.

In one embodiment, the transition metal doped shell is between 2 nm and100 nm thick. In one embodiment, the transition metal doped shell isbetween 2 nm and 80 nm thick. In one embodiment, the transition metaldoped shell is between 2 nm and 60 nm thick. In one embodiment, thetransition metal doped shell is between 2 nm and 40 nm thick. In oneembodiment, the transition metal doped shell is between 2 nm and 20 nmthick. In one embodiment, the transition metal doped shell is between 2nm and 15 nm thick. In one embodiment, the transition metal doped shellis between 3 nm and 14 nm thick. In one embodiment, the transition metaldoped shell is between 4 nm and 13 nm thick. In one embodiment, thetransition metal doped shell is between 5 nm and 12 nm thick. In oneembodiment, the transition metal doped shell is about 5 nm thick. In oneembodiment, the transition metal doped shell is about 10 nm thick.

The transition metal doped shell can be deposited using any method knownin the art, including, but not limited to, thin film sol-gel chemistry,atomic layer deposition, chemical vapor deposition, and sputtering.

In one embodiment, the transition metal doped shell comprises at leastone layer. In one embodiment, the transition metal doped shell comprisesone layer. In one embodiment, the transition metal doped core comprisestwo layers.

In one embodiment, the surface of the transition metal doped shellcomprises a functionalizable group. Exemplary functionalizable groupsinclude hydroxide groups (—OH), amine groups (—NH₂), silane groups(—SiH₃), siloxane groups (—OSiR₃) and thiol groups (—SH). In oneembodiment, the surface of the transition metal doped shell comprisesfree hydroxide groups.

In one embodiment, the surface of the transition metal doped shellexhibits a surface dipole. In one embodiment, the surface dipoleinteracts with the energy levels of the transition metal in thetransition metal doped shell. In one embodiment, tuning the surfacedipole can tune the energy levels in the transition metal in thetransition metal doped shell. In one embodiment, the surface dipole canbe tuned by functionalizing the surface functionalizable groups, such asthrough chemical reaction or through changes in the chemicalenvironment.

In one embodiment, the transition metal in the transition metal dopedshell facilitates upconversion, wherein the light emitted by thecomposite material is higher in energy than the light absorbed by thecomposite material. In one embodiment, the transition metal in thetransition metal doped core facilitates downconversion, wherein thelight emitted by the composite material is lower in energy than thelight absorbed by the composite material. In one embodiment, thetransition metal in the transition metal doped shell sensitizes the rareearth in the rare earth doped core.

In one embodiment, functionalization of the surface groups on thesurface of the transition metal doped shell changes the absorbanceand/or emission energies of the resulting composite material. In oneembodiment, functionalization of the surface groups on the surface ofthe transition metal doped shell decreases the emission wavelength ofthe composite material relative to a non-functionalized compositematerial. In one embodiment, functionalization of the surface groups onthe surface of the transition metal doped shell increases the emissionwavelength of the composite material relative to a non-functionalizedcomposite material.

In one embodiment, the functionalizable groups on the surface of thetransition metal doped shell are treated with at least one carboxylicacid. In one embodiment, the functionalizable groups on the surface ofthe transition metal are treated with a mixture of carboxylic acids.

Inks

In one aspect, the present invention relates to inks comprising thecomposite material of the instant invention. In one embodiment, the inkcomprises glycerol. In one embodiment, the ink comprises at least onecomposite material. In some embodiments, the ink comprises more than onecomposite material. In one embodiment, the composite material in the inkpossesses surface functionalizable groups as described elsewhere herein.

The inks may be used for any purpose known to one of skill in the art.In one embodiment, ink comprising the composite material of the instantinvention is used in the printing of currency. In one embodiment, inkcomprising the composite material of the instant invention is used inthe manufacture of clothing. In one embodiment, ink comprising thecomposite material of the instant invention is used in the printing ofdocuments.

In one embodiment, ink comprising the composite material of the instantinvention is used in the printing of quick response “QR” codes. In oneembodiment, the ink comprising the composite material of the instantinvention further comprises at least one dye. In one embodiment, inkcomprising the composite material of the instant invention presents adifferent “QR” code upon exposure to a light source such as UV, visible,or infrared light.

Methods

In one aspect, the present invention relates to a method of tuning theabsorbance/emission spectrum of a composite material by functionalizingfunctionalizable groups on the surface of the composite material.

In one embodiment, the surface dipole of the surface of the transitionmetal doped shell is tuned by chemical modification of thefunctionalizable groups. In one embodiment, the chemical modificationcomprises a self-limiting carboxylic reaction. In one embodiment,treating the functionalizable group on the surface of the transitionmetal doped shell with a carboxylic acid changes the surface dipole ofthe surface of the transition metal doped shell. In one embodiment,changing the surface dipole of the surface of the transition metal dopedshell changes the coordination geometry of the transition metal element.In one embodiment, changing the surface dipole of the surface of thetransition metal doped shell changes the absorbance or emission spectrumof the transition metal element. In one embodiment, different carboxylicacids produce different changes in absorbance or emission spectra.

Exemplary carboxylic acids include, but are not limited to, formic acid;alkyl carboxylic acids such as acetic acid, propionic acid, and butyricacid; halo-alkyl acids such as chloroacetic acid, dichloroacetic acid,trichloroacetic acid, and trifluoroacetic acid; aromatic carboxylicacids such as benzoic acid, para-(fluorosulfonyl)benzoic acid,para-nitrobenzoic acid, para-cyanobenzoic acid,para-(trifluoromethyl)benzoic acid, para-methoxybenzoic acid,para-bromobenzoic acid, para-chlorobenzoic acid, para-fluorobenzoicacid, para-aminobenzoic acid, para-mercaptobenzoic acid,benzene-1,4-dicarboxylic acid, benzene-1,3-dicarboxylic acid, andbenzene-1,2-dicarboxylic acid; amino acids such as biotin; and diacidssuch as oxalic acid, malonic acid, succinic acid, maleic acid, fumaricacid, itaconic acid, and tartaric acid; organic dyes such as 5(6)-FAM,5(6)-TAMRA, 5(6)-carboxyfluorescein, 5(6)-carboxynaphthofluorescein,5-FAM, 5-ROX, 5-TAMRA, 6-FAM, 6-ROX, 6-TAMRA,7-diethylaminocoumarin-3-carboxylic acid,7-hydroxy-4-methylcoumarin-3-acetic acid, 7-hydroxycoumarin-3-carboxylicacid, 7-methoxycoumarin-3-carboxylic acid, BODIPY® FL, BODIPY® FL C5,DMACA, NBD-X, Oregon Green® 488 carboxylic acid (5-isomer), OregonGreen® 514 carboxylic acid, carboxymethylthiobimane,5-(and-6)-carboxyl-2′,7′-dichlorofluorescein diacetate, Cy3 carboxylicacid, monosulfo Cy3 carboxylic acid, disulfo Cy3 carboxylic acid, Cy5carboxylic acid, monosulfo Cy5 carboxylic acid, disulfo Cy5 carboxylicacid, Cy5.5 carboxylic acid, Cy7 carboxylic acid, and disulfo Cy7carboxylic acid; solubility enhancers such as polyethylene glycolcarboxylic acid (mPEG-COOH), O-(2-carboxyethyl)-O′-methyl-undecaethyleneglycol, 2-[2-(2-methoxyethoxy)ethoxy]acetic acid, methoxypolyethyleneglycol acetic acid, methoxypolyethylene glycol propionic acid,O-methyl-O′-succinylpolyethylene glycol, andO-[2-(3-succinoylamino)ethyl]-O′-methyl-polyethylene glycol. In oneembodiment, the carboxylic acid is a para-substituted benzoic acid. Inone embodiment, the carboxylic acid is selected from the groupconsisting of para-(fluorosulfonyl)benzoic acid, para-nitrobenzoic acid,para-cyanobenzoic acid, para-bromobenzoic acid, benzoic acid,para-methoxybenzoic acid, and para-aminobenzoic acid. In one embodiment,the carboxylic acid is para-nitrobenzoic acid. In one embodiment, thecarboxylic acid is para-aminobenzoic acid.

In another aspect, the present invention relates to a method ofidentifying counterfeit currency. In one embodiment, the method includesthe steps of providing an ink having a core-shell nanoparticle withsurface functionalizable groups; applying the ink during the minting ofauthentic currency; treating a currency sample with a solutioncomprising a carboxylic acid; and exposing the currency sample to UVlight. In one embodiment, the core-shell nanoparticle comprises atransition metal doped shell disposed over a rare earth doped core.

In one embodiment, the method can be used to identify counterfeitcurrency. In one embodiment, the currency is a paper currency. In oneembodiment, the currency is minted, such as by a governing body. In oneembodiment, at least one region of the currency is printed using an inkcomprising the inventive composite material. In one embodiment, at leastone region of the currency is printed using an ink comprising acore-shell nanoparticle with surface functionalizable groups. In oneembodiment, an ink comprising a core-shell nanoparticle with surfacefunctionalizable groups is applied during the minting process.

In one embodiment, a currency sample is provided. The currency isnecessarily either authentic currency comprising a core-shellnanoparticle with surface functionalizable groups or counterfeitcurrency lacking a core-shell nanoparticle with surface functionalizablegroups.

In one embodiment, the currency sample is exposed to UV light. In oneembodiment, the core-shell nanoparticle with surface functionalizablegroups emits a specific wavelength of visible light upon exposure to UVlight. In one embodiment, counterfeit currency (currency not comprisinga core-shell nanoparticle having surface functionalizable groups) doesnot emit the same wavelength of light upon exposure to UV light.

In one embodiment, the currency sample is treated with a carboxylic acidsolution. In one embodiment, the carboxylic acid solution reacts withthe core-shell nanoparticle with surface functionalizable groups andthereby changes the emission wavelength of the core-shell nanoparticleupon exposure to UV light.

In one embodiment, the currency treated with a carboxylic acid solutionis exposed to UV light. In one embodiment, upon exposure to UV light,currency comprising a core-shell nanoparticle emits a specificwavelength of visible light that is visibly distinct from the wavelengthemitted upon UV light exposure prior to the carboxylic acid treatment.In one embodiment, counterfeit currency (currency not comprising acore-shell nanoparticle having surface functionalizable groups) does notemit the same wavelength of light upon treatment with a carboxylic acidand subsequent exposure to UV light. In one embodiment, the lack ofvisible light emission distinguishes counterfeit currency from authenticcurrency.

EXPERIMENTAL EXAMPLES

The invention is further described in detail by reference to thefollowing experimental examples. These examples are provided forpurposes of illustration only, and are not intended to be limitingunless otherwise specified. Thus, the invention should in no way beconstrued as being limited to the following examples, but rather, shouldbe construed to encompass any and all variations which become evident asa result of the teaching provided herein.

Without further description, it is believed that one of ordinary skillin the art can, using the preceding description and the followingillustrative examples, make and utilize the composite materials of thepresent invention and practice the claimed methods. The followingworking examples therefore, specifically point out the preferredembodiments of the present invention, and are not to be construed aslimiting in any way the remainder of the disclosure.

Example 1: Tunable Nanophosphors

TM ions are doped into a ˜10 nm TiO₂ shell (Wei Seh, et al., Nat.Commun. 2013, 4, 1331) surrounding a 30 nm RE-doped 3-NaYF₄ crystal(Zhengquan and Yong, Nanotechnology 2008, 19, 345606). An exemplarynanoparticle is shown in FIG. 1. These two different host lattices allowfor chemical stability while simultaneously protecting the luminescentcenter for parasitic phonon vibrations (Suyver, et al., J. Lumin. 2006,117, 1-12; Dorman, et al., J. Phys. Chem. C 2014, 118, 16672-16679). Twodifferent sets of dopant combinations are investigated, one fordownconversion and one for upconversion luminescent tuning. Thedownconversion pair consists of Tb³⁺ and Cr³⁺, in order to takeadvantage of the strong green emission with a tunable green absorption,respectively, as shown in FIG. 2 (Morosin, Acta Crystallogr. Sect. B:Struct. Crystallogr. Cryst. Chem. 1972, 28, 1899-1903; Dorman, et al.,J. Phys. Chem. C 2012, 116, 12854-12860). Upconversion ion pairs consistof Er³⁺ with Ni²⁺ or Cu²⁺. In this case, the TM ion was chosen toamplify the light absorption and energy transfer to Er³⁺, which has beenshown to produce a blue, green, or red emission, depending on the dopantconcentration. By changing the absorption spectrum of the TM ion, themagnitude of energy transfer, and the number of excitation steps neededfor a specific transition, can be controlled (FIG. 2).

Dipole moments are then applied to the surface using a self-limitingcarboxylic reaction in which surface hydroxide groups are replaced withvarious benzoic acid derivatives (Table 1) (Goh, et al., J. Appl. Phys.2007, 101, 114503-114503). These groups have a difference in dipolemoments up to 9 Debye, slightly weaker than the highly ionic KBr salt(10.5 Debye). After the luminescent range is quantified, thenanoparticles are printed using an inkjet printer. This process allowsfor plenty of optimization in order to maintain desired luminescentproperties. Specifically, both aqueous and organic solvents can be usedto print these phosphors, but both require the use a glycerol as astabilizing agent (You, et al., Nanoscale 2015, 7, 4423-4431). While theglycerol poses no major issues, the optimization and characterization ofthis process is performed in-depth in order to quantify the role of thisstabilizing agent. Standard characterization is performed to quantifyphosphor optical performance including standard and transientabsorption, excitation, and emission spectra. The core-shellnanoarchitecture is verified through standard crystal characterizationtechniques to elucidate crystal structures and compositions at a sharedinstrument facility, which is employed in the characterization of TM-REinteractions, local binding environment, and energy level landscapes.

TABLE 1 Functional Dipole Group (-R) Moment (D) SO₂F 4.5 NO₂ 3.8 CN 3.4Br 1.4 H −2.1 OCH₃ −3.9 NH₂ −4.5

Data collection and analysis is performed using the standard techniquesfor luminescent phosphors, including UV-Vis, Photoluminescence,Transient absorption spectroscopy, and luminescent decay kinetics(Dorman, et al., J. Phys. Chem. C 2012, 116, 10333-10340; Dorman, etal., J. Phys. Chem. C 2012, 116, 12854-12860; Dorman, et al., J. Appl.Phys. 2012, 111, 083529), structural characterization (TEM, XRD, SEM),and chemical characterization (XPS, EDX). Nonlinear spectroscopy such assecond harmonic generation, upconversion spectroscopy, and two-photonfluorescence spectroscopy is coupled with ultrafast transient absorptionspectroscopy to carefully study the photodynamics of the preparednanomaterials (Karam, et al., J. Chem. Phys. 2016, 144, 124704; Kumal,et al., Langmuir 2015, 31, 9983-9990). The optimization of the inkjetprinting process is based on the wettability of the solution on variousbrands and types of paper for direct application. This process alsoincludes precise control of solution viscosities and particleconcentrations such that the doped ink is printed at 2,000 dpi, aspacing of roughly 10 Lm between phosphors. This resolution is possiblebased on currently available inkjet printer's technical specifications(Epson Stylus Photo R3000 Inkjet Printer).

Example 2: Synthesis, Modification, and Optical Properties of DopedMetal Oxides

In order to produce high-efficiency luminescence, the structure andpurity of the crystal are important, requiring carefully designedsynthetic methods. For RE-based luminescence, yttrium based compoundsare typically employed to allow for RE dopant incorporation, reducingthe number of defect states due to charge compensation (Zhou, et al.,Chemical Reviews 2015, 115, 395-465). In order to avoid oxygen defects,high-quality nanostructures are preferred. For example, yttriumnanostructures have been fabricated using the hydrothermal method,resulting in “lossy” Y(OH)₄ instead of the preferred Y₂O₃ (Dorman, etal., The Journal of Physical Chemistry C 2010, 114, 17422-17427). TheY(OH)₄ structures were converted using high-temperature annealing.Unfortunately, annealing allows the diffusion of dopant atoms,necessitating a detailed understanding of the conversion steps (FIG.3A). By probing the bond length transformations via in situ XAS, therequired annealing temperature and time was extracted to limit dopantaggregation. However, the overlap in RE energy levels produces unwantedenergy transfer between ions.

Core-shell nanostructures, produced through sol-gel modification,facilitate the segregation of ions, with ±2 nm precision (Dorman, etal., The Journal of Physical Chemistry C 2012, 116, 10333-10340). Thistechnique produced sharp interfaces as highlighted by the Y₂O₃|Yb₂O₃interface (FIG. 3B). These techniques were employed to develop acore-multi shell structure to promote and hinder energy transfer betweenRE dopants for high-quality white light (Dorman, et al., The Journal ofPhysical Chemistry C 2012, 116, 12854-12860). The quality of the whitelight was engineered through dopant position and shell layer thicknessto produce a “natural” soft spectrum (FIG. 3C) for commercial LEDs. Thecombination of the four different RE elements also resulted in a highcolor rendering index, outlined by the polygon, typically not possiblewith two and three wavelength LEDs. Finally, similar metal oxideparticles were modified using the self-assembly of carboxylic acidmolecules to attach to surface hydroxides over 2+ hrs, binding dyes tothe surface (not shown) (Dorman, et al., The Journal of PhysicalChemistry C 2014, 118, 16672-16679; Weickert, et al., APL Materials2013, 1, 042109). This technique is also valid for dipoles (Table 1) andbiological groups.

Ni²⁺ doped metal oxide thin films and core-shell structures have beensynthesized and characterized. Ni²⁺ was selected as an initial TM dopantdue to its strong blue/red absorption (Wenger, et al., Journal of theAmerican Chemical Society 2000, 122, 7408-7409). The core and valencelevels have been probed using soft XAS (CAMD), x-ray photoemissionspectroscopy (XPS), and UV-Vis absorption (FIG. 4). Thin film depositionuses a modified TiO₂ sol-gel chemistry where titanium isopropoxide isaged 24 hrs in an acidic ethanol solution and spin coated on glass andSi wafers (Yu, et al., Scientific Reports 2015, 5, 9561). Thin filmswere functionalized with two para-benzoic acid (BZA) ligands, one withan NO₂ group (electron withdrawing, 3.8 D, c.f. Table 1) and one with anNH₂ group (electron donating, −4.5 D), to control the electron densityat the interface.

In order to study both sides of the Ni 3d-O ²p hybridization, O K edge(FIG. 4A) and Ni LIII/LII edges (FIG. 4B) were probed via XAS. Eachregion was scanned with a 0.1 eV step size and absorption energiescalibrated to reference NiO or TiO₂ films. The post O K edge (537-550eV) scans show the effect of the ligand on the Ni—O hybridization withenergy shifts proportional to the dipole moment, with peak maxima of theNH₂-BZA modified surface 0.5 eV high than NiO and the NO₂-BZA modifiedsurface 1 eV lower than TiO₂ references. This dipole effect is also seenin the LII edge of the Ni absorption, where the e_(g) edge increases inenergy as NO₂-BZA (882.3 eV) is replaced with NH₂-BZA (882.5 eV) and thet_(2g)/e_(g) peak ratio changes from 0.99 to 1.01. These shifts areattributed to modification of the orbital splitting (Δ), electronfilling, and spin-orbital coupling (Ogasawara, et al., Physical Review B2001, 64, 115413). Additionally, XPS scans (FIG. 4C) show the ability toaffect the intensity of the core Ni 3p level with surface dipoles.Conversion from NO₂ to NH₂ groups decrease intensity by ˜20% in relationto the normalized Ti 3s peak. Furthermore, it was important todemonstrate the effect of these dipoles on the optical transitions. Thindoped and pure Y(OH)₃ films (Dorman, et al., The Journal of PhysicalChemistry C 2012, 116, 10333-10340) were deposited around ˜500 nmhexagonal 3-NaYF₄:Er³⁺ nanoparticles to measure Ni absorption (Li, etal., Inorganic Chemistry 2007, 46, 6329-6337). FIG. 4D shows theabsorption of the Er³⁺ doped core (+), the Ni²⁺ doped shell, and thefunctionalized core-shell structures. Upon functionalization withNH₂-BZA, the characteristic Ni absorption (˜410 nm) red shifts up to 20nm (broad absorption peak centered at 435 nm). This peak disappearscompletely with the NO₂-BZA group, likely moving into the Y(OH)₃/NaYF₄band gap. These results demonstrate the ability to quantify crystalfield manipulation with optical and electronic measurements.

These materials are modeled to extract the electronic structures usingTD-DFT (Govind, et al., The Journal of Physical Chemistry Letters 2011,2, 2696-2701). Simulations of the TiO₂ O K edge were performed using a 7atom anatase TiO₂ cluster, passivated with hydrogen atoms, to findappropriate basis sets. The preliminary basis set and cluster(Aug-CC-pvTZ, FIG. 5 may require further optimization though exactmatching is not expected, as this small cluster cannot account forspin-orbital splitting or inhomogeneity in the crystal structure. Thehigh energy peaks, centered around 540 eV, can be attributed to discretelevels in the model extending beyond the vacuum level (Govind, et al.,The Journal of Physical Chemistry Letters 2011, 2, 2696-2701).

Example 3: Atomically Controlled Nanostructures for Directed d-OrbitalModification

Wet chemical methods have been developed to produce metal oxides andfluorides with a wide range of morphologies and sizes (Dorman, et al.,The Journal of Physical Chemistry C 2014, 118, 16672-16679; Dorman, etal., The Journal of Physical Chemistry C 2012, 116, 10333-10340; Dorman,et al., The Journal of Physical Chemistry C 2012, 116, 12854-12860;Dorman, et al., The Journal of Physical Chemistry C 2010, 114,17422-17427; Li, et al., Inorganic Chemistry 2007, 46, 6329-6337; Wang,et al., Chemistry of Materials 2007, 19, 727-734). The ability tomanipulate nanostructures is used to probe uniform and non-uniformligand fields and their effect on TM optical properties. As such,NaYF₄₁TiO₂:Ni²⁺ core-shell structures are synthesized and characterizedand the resulting optical and electronic properties measured with andwithout surface functionalization. The synthesis concentrates on 0D and1D morphologies with critical dimensions around 20-30 nm to account forisotropic (0D) and anisotropic (1D) ligand fields. These nanostructuresare synthesized using hydrothermal/solvothermal (Orgel, J. Chem. Soc.1952, 4756-4761; Yuan, et al., Journal of the American Chemical Society2013, 135, 8842-8845; Wang, et al., Chemistry of Materials 2007, 19,727-734), molten salt (Dorman, et al., The Journal of Physical ChemistryC 2012, 116, 10333-10340; Dorman, et al., Journal of Applied Physics2012, 111, 083529), and thermal decomposition methods with controlleddimensionalities. For example, hexagonal NaYF₄ nanorods/wires have beensynthesized hydrothermally by heating an acidic solution of Y(NO₃)₃ andsodium citrate at 180° C. Crystallite diameters, lengths, and aspectratios can be tuned via the RE:Citrate precursor ratio (Li, et al.,Inorganic Chemistry 2007, 46, 6329-6337). Alternatively, 25 nm hexagonalβ-NaYF₄ can be synthesized via thermal decomposition of YCl₃ salts in anon-aqueous solution containing NH₄F at temperatures up to 350° C.(Ostrowski, et al., ACS Nano 2012, 6, 2686-2692).

Thin shell layers (˜5 nm) of TiO₂:Ni²⁺ metal oxides are deposited usinga similar thin film sol-gel chemistry (Li and Zhao, Advanced Materials2013, 25, 142-149; Lu, et al., Catalysis Science & Technology 2016, 6,6845-6852; Tang, et al., ACS Catalysis 2013, 3, 405-412), to act as thescaffolding layer for further functionalization. Additionally, Y₂O₃(Dorman, et al., The Journal of Physical Chemistry C 2012, 116,10333-10340) and SiO₂ layer depositions serve as alternatives. Finally,the particle surface ware functionalized with multiple dipole groups(Table 1). These experiments focus on established NO₂-BZA and NH₂-BZAligands, spanning 8.3 D, to measure the full Ni²⁺ response. In additionto these two groups, the table includes other benzoic acid derivatives,biomolecules, and linking groups which can be employed in the future forhigher sensitivities and inter/intra-particle energy transferexperiments. Atomic location (STEM/EELS) and the electronic effects ofligands (FIG. 4 c) are quantified and coupled with UV-Vis. STEM/EELSanalysis has been performed on core-shell Fe₃O₄|MnFe₂O₄ nanoparticles,as seen in FIG. 6, and shows the location of Fe (green) and Mn (purple)ions. These methods demonstrate control of synthetic methods and abilityto modify electronic and optical properties. Finally, surface coverageis exploited to further modify ligand field strengths.

During functionalization, the movement of electrons has both electronicand structural consequences which are not seen under standard methods.As previously discussed, there is evidence of evolving Ni—O and Ti—Obond structures with surface functionalization. In order to investigatethese effects, high-resolution, in situ characterization on theoptimized 0D/1D particles is performed during surface functionalization.Additionally, surface adsorption measurements are coupled with opticalcharacterization equipment to calibrate the surface functionalizationkinetics and absorption/luminescence shifts.

Prior to synchrotron measurements, adsorption kinetics are investigatedvia light absorption (Lambda 950) and photoluminescence (Edinburgh FLS980 Spectrofluorometer with time correlated single photon counting(TCSPC)). The two systems allow for independent but related measurementsof the shifting electronic landscape during ligand bonding. Thestructural and electronic properties of these materials are investigatedusing two complementary methods, XAS and UPS, to probe the shift invalence and unoccupied energy levels, and structural modifications. Nearedge XAS (VLSPGM beamline, 200-1200 eV) is used to extract the valencestate, coordination chemistry, and electronic structure of theunoccupied states. When applicable (>1 keV between absorption edges) theextended region (EXAFS) is collected to further elucidate the dopantcoordination chemistry. Additionally, the electronic state of thevalence and outer core levels, up to 75 eV binding energies, is probedvia UPS (5m-TGM beamline, 25-240 eV). Due to the low energies beingemployed, the chamber(s) require low pressures to avoid x-rayabsorption. Dipole surface coverage was increased incrementally betweenscans by controlling the partial pressure, via heating of the powders,and exposure times (Lu, et al., Catalysis Science & Technology 2016, 6,6845-6852).

Localized external fields can be used to manipulate d-orbitalhybridization, i.e., internal crystal fields. The measure-model aspectof this work is achieved by comparing crystal field parameters ofmeasured and simulated 0D/1D core-shell structures. Crystal fieldstrengths can be extracted from optical measurements in the form of theRacah crystal field parameter (B) using Tanabe-Sugano diagrams as idealsystems. Both the crystal field splitting (Δ) and Racah parameters havebeen linked to M-O bond distance (Bocquet, et al., Physical Review B1992, 46, 3771-3784; Hauser, In: Spin Crossover in Transition MetalCompounds I, Gitlich, P.; Goodwin, H. A., Eds. Springer BerlinHeidelberg: Berlin, Heidelberg, 2004; pp 49-58). Under an octahedralsite symmetry assumption, the crystal field strength is equal in alldirections (0D structure). As field anisotropy increases, the octahedralassumption is no longer valid, as the absorption spectrum no longermatches the octahedral model. One example of this process is tetragonalelongation (i.e., expansion along the z-axis) which causes a contractionof the d_(xy) and d_(z) ₂ orbitals (Burdett, Inorganic Chemistry 1981,20, 1959-1962). While experimental data offers a solid base for theprediction of crystal field and expected physical properties,correlating these results with simulation can provide furtherjustification for the observed effects.

Simulations have been performed using TD-DFT. While this technique iscommonly used to model optical and electronic properties of molecularcompounds, it can also be applied to solid structures (De Angelis, etal., Nanotechnology 2008, 19, 424002). In order to increase the accuracyof the model, a 2×5×2 bulk-mimicking cluster is employed to modelTiO₂:Ni²⁺ for TD-DFT simulations (FIG. 7A). The simulation enables theattachment of ligands to multiple surface facets by cleaving thestructure after optimization (FIG. 7B). The simulations are used tovalidate the experimental data via substitution of hydrated surfaceswith functionalities possessing various dipole moments. The TD-DFTresults can then be treated similarly to the measured data fordetermination of B. Additionally, this process is used to produceanisotropic and isotropic fields to model d-orbitalcontraction/elongation. The effect of field anisotropy is hypothesizedto be an important factor in d-orbital hybridization.

Despite their susceptible nature, TM dopants are attractive alternativesto RE dopants due to their high absorption cross-section. This has ledto the incorporation of Mn²⁺ (Dan, et al., Materials Letters 2015, 150,76-80; Tian, et al., Advanced Materials 2012, 24, 1226-1231) or Ni²⁺(Takeda, et al., Journal of Applied Physics 2016, 120, 073102) assensitizers to increase phosphorescence or upconversion yield. However,these studies are restricted to specific wavelengths where there waslittle variability in the dopant absorption, such as the ³A₂→¹Etransition in octahedral Ni²⁺, which absorbs solely at 1100 nm (Takeda,et al., Journal of Applied Physics 2016, 120, 073102) (FIG. 8). Whilethe ultimate goal in luminescence is to replace RE dopants, this can beviewed as a two-step process, first replacing the high concentrationsensitizer ions (Yb³⁺/Ce³⁺, up to 20 mol %) then the emitter ion. Inaddition to energy level matching, RE dopants demonstrate bothdownconversion (UV→visible) and upconversion (IR→visible) luminescence.Therefore, the TM energy levels can be tuned to enhance or quenchspecific emission processes.

The ability to direct energy transfer between various layers withincore-shell structures for enhanced luminescent properties has beenpreviously demonstrated. One such RE ion pair that demonstrates energytransfer in both luminescence mechanisms is the Er³⁺—Yb³⁺ system(Dorman, et al., The Journal of Physical Chemistry C 2012, 116,10333-10340; Rodriguez, et al., Solar Energy Materials and Solar Cells2010, 94, 1612-1617). This system is widely used, with applications infiber optics, prompting this investigation to remove the highconcentration Yb sensitizer (20 mol %) (Agrawal, G. P., Nonlinear FiberOptics; Academic press, 2007). Accordingly, the focus is on dictatingthe Ni²⁺—Er³⁺ coupling, and energy transfer, by tuning the TM excitedstate energies to match those of the RE ions. The Ni²⁺—Er³⁺ system isshown in FIG. 8 with expected coupling regions shown in the dashedlines. Incorporation of Er³⁺ into the core can be readily performed viasubstitution of yttrium precursors with the Er analog (Dorman, et al.,The Journal of Physical Chemistry C 2012, 116, 12854-12860; Li, et al.,Inorganic Chemistry 2007, 46, 6329-6337). Initial coupling experimentsare performed on downconversion processes to turn emission on/off(energy sink) or vary excitation/emission pathways (coupling mechanism).This process is preferred due to the higher RE emission efficiencies,their ability to excite without a laser source, and the wide range ofavailable energies for the TM dopants (FIG. 8).

The Ni²⁺—Er³⁺ coupling is probed using an Edinburgh FLS 980spectrofluorometer. This system is equipped with dual monochromators toallow for the simultaneous screening of excitation and emissionwavelengths. The equipment facilitates dynamic studies which are used toextract surface functionalization dependent luminescence and excitedstate Ni²⁺—Er³⁺ energy transfer processes. Additionally, time-resolvedluminescence mechanisms are extracted using TCSPC to quantifyefficiencies (Dorman, et al., The Journal of Physical Chemistry C 2012,116, 10333-10340). Low-temperature luminescence removes thermalquenching caused by the host materials and can result in discrete TMexcitation energies. These methods have been employed to extract excitedstate kinetic processes for RE-RE transitions; here, this analysis isextended to the RE-TM system in order to understand and optimize thecoupling between the dopant types (Broholm, C.; Fisher, I.; Moore, J.;Murnane, M., Basic Research Needs Workshop on Quantum Materials forEnergy Relevant Technology. Energy, D. o., Ed. 2017).

The culmination of this work is the establishment of a set of designrules to describe the electro-optical coupling effects at reduceddimensions. As dimensionality decreases, a new method of predictingorbital energies and expected physical properties is needed. Bydeveloping a fundamental understanding of the effects of localizedfields on symmetry breaking (hybridized d-orbitals), new optical,magnetic, and catalytic materials can be designed. Theoretical modelshave been employed to help optimize the coupling between crystalstructures/bonds and optical properties and to improve luminescentefficiencies (Dorman, et al., Journal of Applied Physics 2012, 111,083529). Using the electronic, structural, optical, and theoreticalresults from the previous objectives, design rules are formulated. Theserules are then tested on alternative RE-TM systems to determine theiraccuracy as a universal prediction tool. The developed model is appliedto upconversion luminescence in an effort to increase luminescentefficiencies while simultaneously reducing the concentration of REdopants. The criteria revolve around the propensity for orbitalmanipulation, a range of possible crystal field strengths, andobserved/allowed optical transitions. Data are collected in a databaseand distributed for broader integration into existing communities (e.g.,LBNL's Materials Project).

Highlighting the ability to control the different energy levels of theNi in the TiO₂ phosphor allows the control of optical properties and howenergy is transferred to rare earths (FIG. 9A and FIG. 9B). There areslight changes in the electronic structure of the host which can beemployed in solar cells and other optoelectronic applications where theelectronic structure plays a role in device performance (FIG. 10). Theseparated Ni L edge experimental and modeling results are shown in FIG.11A and FIG. 11B, respectively, and suggest that the Ni atoms adopt anelongated octahedron geometry when the surface is functionalized withpara-benzoic acid ligands. The t_(2g):e_(g) intensity ratio inversionfor NO₂-BZA bonded films can only be modeled by an elongated octahedron(FIG. 11B, insert), which is done by setting the same orbital hopingparameters as 1.8. Preliminary results suggest that it may be possibleto model the electronic state of the TiO₂ and Ni doped TiO₂ structuresso that expected results can be coupled to experimental results (FIG.12).

Importantly, the absorption can be tuned to the degree of up to 70 nmusing the same host material; this structure can be used to sensitizeblue and green upconversion emission (FIG. 13). This is necessary forcounterfeit technologies to tune the appearance. Er and Ni are coupledin a core-shell structure to enhance energy transfer. Ni has been shownto enhance luminescent efficiency of Er based phosphors.

Careful selection of processing conditions prevents the segregation ofdopants (here, Ni) within the final phosphor (FIG. 14A). Success inpreventing segregation can be assessed by comparing optical propertiesto an NiO reference (FIG. 14B). Initial results suggest that the Nisegregates out based on the amount of moisture it has been exposed to.The amount of moisture can be tracked using infrared spectroscopy (FIG.15A) and the light absorption or color (FIG. 15B) of the pre-annealedsamples. If annealed at the correct point, the Ni does not leave theTiO₂ lattice and the optical and electronic properties of the particlesare fixed for solar/photocatalytic applications. Undesired dopantsegregation is not unique to Ni. The same segregation was observed inother first row transition metals, such as Co as seen in crystalstructure (FIG. 16A) and optical properties (FIG. 16B). The resultssuggest that annealing parameters represent a further variable that canbe manipulated to tune the light absorption for adaptive luminescenceapplications.

Some dopants, such as iron (FIG. 17A) or chromium (FIG. 17B), change thestructure of the TiO₂ to rutile upon annealing, but segregation is lesslikely. These doped species still have very prominent light absorptionand colors. The differing behavior of transition metal dopants withregard to segregation may be due in part to differences in ionic radii.The previously observed phenomenon of dopant segregation with Ni²⁺ andCo²⁺ is likely due to their larger ionic radii compared to that of hostcation (Ti⁴⁺). In the case of Fe²⁺ and Cr³⁺, the annealing rate has noeffect on segregation of dopant metal oxide clusters. This differencemay be attributed in part to the incorporation of smaller sized dopantions in the host lattice of TiO₂.

Weak fields can be used to tune the excitation wavelengths in YVO₄:Eu³⁺(5 mol %)|YVO₄:Bi³⁺ (6 mol %) core-shell NPs similar to the adaptiveabsorption observed in the NaYF₄:Er|Y(OH)₃:Ni core-shell NPs. Bi is acommon rare earth sensitizer but specifically works in YVO₄. Thecore-shell structure facilitates the tuning of emission intensity andexcitation wavelengths so that light appearance may be directlycontrolled. Traditionally, these compounds are used in biologicalimaging. FIGS. 18A and 18B show how a ligand/biomolecule can be attachedto the surface to enhance or quench the luminescence. The resultingnanoparticle represents a novel low-cost biological sensor which canbind to specific biological species for rapid detection. A slightdecrease in luminescent lifetime is observed upon with ligand attachment(FIG. 19A and FIG. 19B). This decrease is expected and is proportionalto luminescent efficiency. Typically, these structures would result inmore prominent reduction in lifetimes and decreases in efficiency.

Example 4: Weak Field Tuning of Transition-Metal Dopant Hybridization inSolid Hosts

Previous reports focusing on tailoring the hybridization in a crystalare based on the application of strong electric/magnetic fields(Shanavas, et al., Phys. Rev. Lett. 2014, 112, 086802; Asamitsu, et al.,Nature 1995, 373, 407), mechanical stress (Chen, et al., Appl. Phys.Lett. 2011, 98, 241916), or crystal composition (Dai, et al., J. Mater.Chem. C 2013, 1, 4570-4576; Bailey, et al., J. Am. Chem. Soc. 2003, 125,7100-7106) to distort local symmetries or crystal field splitting.Strain engineering has been commonly implemented to indirectly controlthe metal-oxygen hybridization by changing the Ni—O bond distance inNi-doped SrTiO₃ thin films (Bai, et al., Sci. Rep. 2014, 4, 5724) andorgano-metallic complexes (Byrne, et al., Chem.-Eur. J. 2012, 18,7738-7748). However, these methods are limited due to the constantapplication of strong fields or the irreversible modification of thecomposition, which can limit device performance via space chargesaccumulation or unwanted geometric distortion (Skobel'

yn, D. V., Surface Properties of Semiconductors and Dynamics of IonicCrystals; Plenum Press: New York, U.S., 1971; Vol. 48). Therefore, itremains a challenge for the scientific community to modify thehybridization of atomic orbitals in a stable but reversible manner.

Reversible tuning of crystal field splitting energy for controlling theTM 3d-O 2p hybridization in a TM-doped solid is possible using weakexternal fields (surface dipoles) (Goh, et al., Appl. Phys. 2007, 101,114503). While the overall effect of these external fields is limitedprimarily at the surface (Δ₀ ∝ R⁻⁵, where R is the metal-ligand distancein octahedral TM complexes) (Moreno, et al., Int. J. Quantum Chem. 1994,52, 829-835) they can be reversibly manipulated (Lopez-Sanchez, J. A.,et al., Nat. Chem. 2011, 3, 551). It is possible to manipulate theinterfacial electron density with polarized molecules, potentiallymodifying the TM dopant 3d orbitals/p-d hybridization to control theelectronic properties of a film. This response has been previously usedto tune the open circuit voltage/short circuit current in photovoltaicdevices by systematically controlling the recombination kinetics andcharge injection (Goh, et al., Appl. Phys. 2007, 101, 114503).

To study the relationship between TM-O hybridization in the presence ofsurface dipoles, Ni-doped TiO₂ films were chosen for their chemicalstability and strong optical response. Ni was incorporated into a TiO₂thin film by spin coating a dilute sol-gel solution and annealing. Thethin films were characterized using high-resolution transmissionelectron microscope (HRTEM), with elemental electron loss spectroscopy(EELS) chemical mapping, X-ray Diffraction (XRD), and UV-Visspectroscopy to determine the surface composition, bulk crystalstructure, and crystal field splitting energy of the solid. The surfaceof the TiO₂:Ni films was functionalized with para-substituted benzoicacid ligands which can modify the dipole moment over 8 D. The influenceof these external chemical fields on the electronic structure of theinterfacial dopants is probed via surface sensitive electroniccharacterizations such as X-ray Photoelectron Spectroscopy (XPS) andsoft X-ray Absorption Spectroscopy (XAS). The relationship between thedopant electron density in the valence 3d orbitals and theelectronegativity of the benzoic acid substituent has been studied. Thecharacterization results suggest that the covalency/hybridization of thedopant (Ni)-oxygen (O) bond increases for electronegative substituentsand vice versa for electropositive substituents.

Materials and Methods

Titanium (IV) Isopropoxide (TTIP, Acros Organics, >98%), nickel (II)chloride hexahydrate (NiCl₂.6H₂O, BTC, >99%), hydrochloric acid (HCl,36-38.5% purity, ACS grade), p-nitrobenzoic acid (NO₂-BZA, Alfa Aesar,99%), p-aminobenzoic acid (NH₂-BZA, VWR Chemicals), reagent alcohol(<0.075% VWR Analytical), acetonitrile (HPLC grade, >99.8%) wereobtained commercially. All the materials were used without furtherpurification.

The sol required for coating TiO₂:Ni film was synthesized by employingsol-gel chemistry (Yu, et al., Sci. Rep. 2015, 5, 9561). NiCl₂.6H₂O wasdissolved in 5 mL of ethanol and then 1.5 mL of TTIP was added dropwiseto form homogeneous TiO₂:Ni sol after 3-4 h of continuous stirring. 125μL of HCl was used as a catalyst in this process to control the rapidhydrolysis of TTIP precursor. The concentration of Ni precursor to TTIPwas varied from 0 to 15 mol %. The prepared sol was aged for 24 h beforespin coating. The sol was diluted (1:2, v/v) with ethanol prior to spincoating onto Si(100) substrates at 3000 rpm for 60 s. The spin-coatedsamples were dried at 100° C. for 5 min with subsequent annealing at450° C. for 2.5 h under low vacuum (˜100 mtorr).

The surface of the inorganic film was modified with benzoic acid (BZA)ligands via carboxylic acid chemistry (Goh, et al., Appl. Phys. 2007,101, 114503), wherein the carboxylic groups chemisorb onto thehydrophilic surface. Two para-substituted BZA groups were chosen to actas an electron withdrawing group (NO₂, μ=3.8 D) and an electron donatinggroup (NH₂, μ=−4.5 D). The TiO₂ films were immersed in 1 mM acidsolution in acetonitrile. After 2-3 h, the samples were rinsed withethanol and isopropanol before drying in air.

The thickness of the film was measured using a Filmetrics (F3-UV)reflectometer tool. A standard Si(100) wafer was used as a reference toaccount for the native oxide layer. HRTEM images were obtained(sensitive to light elements) using the 200 kV JEOL-ARM electronmicroscope equipped with double aberration correctors, adual-energy-loss spectrometer, and a cold FEG source. Scanning EELSspectra were obtained with a convergence semi-angle of 20 mrad, and acollection semi-angle of 88 mrad. Dual EELS mode was used to remove theintrinsic energy shifts of the electron beam introduced in the EELSscanning process. The EELS spectra were background subtracted with apower-law function, and multiple scattering was removed by a Fourierdeconvolution method. The elemental maps were determined fromNi-L_(II,III), Ti L_(II,III), O K, and Si L edges. The crystal structurewas identified by performing Gracing Incidence (GI)-XRD usingPANalytical X-ray diffractometer operating at 45 kV and 40 mA. The θ-2θradial scan was performed over the range 15-70° with a step size of0.04° and dwell time of 60 s, using Cu K_(α1) (λ=1.54 Å) as radiationsource.

The absorption spectra of TiO₂:Ni was recorded using a Perkin-ElmerLambda 900 UV/Vis/NIR spectrometer equipped with an integrating sphereand a center-mounted sample holder. The absorption scans ranging from300-1300 nm with a scan rate of 0.5 nm/s were obtained on the thin filmsdeposited on glass substrates before annealing. The change inmonochromators was set to occur at 900 nm. Fourier-transform infrared(FTIR) spectroscopy was performed on the surface functionalized TiO₂thin films using an ATR Germanium crystal in a Thermo Scientific Nicolet6700 FTIR equipped with an MCT detector cooled to liquid N₂temperatures. The data was collected in absorbance (log (1/R)) mode,with air as background, and resolution being 4 cm⁻¹ in the region goingfrom 400-1400 cm⁻¹. The incident angle of the laser was kept at 50° tocollect the total internally reflected light.

The oxidation states of the TiO₂:Ni film were determined from XPSmeasurements performed using Scienta Omicron ESCA 2SR XPS system. Amonochromatic Al K_(α1) X-ray source and a hemispherical analyzer with a128 channel detector were used for all samples. The pressure inside thechamber was maintained at 1.5×10⁻⁹ torr. The XPS spectra were calibratedto adventitious C is peak at 284.6 eV. The step size of thesemeasurements was 0.05 eV and the inherent Gaussian width of the sourcewas 0.167 eV. All peaks were fit (using CasaXPS software) to symmetricVoigt line shapes that were 70% Gaussian and 30% Lorentzian productfunctions.

X-ray Absorption Near Edge Spectroscopy (XANES) measurements were takenat two beamlines: High Energy X-ray Absorption Spectroscopy (HEXAS: 5-30keV) beamline on a 11-pole, 7.5 T multi-pole wiggler and Vacuum LineSpacing-Plane Grating Monochromatic (VLSPGM: 0.2-1 keV) at the Centerfor Advanced Microstructures and Devices (CAMD). Ni K edge in TiO₂:Ni(15 mol %) thin film was measured in the HEXAS beam line in fluorescencemode of detection. The L edges of Ti, Ni and K edge of O were measuredin the VLSPGM beam line with photon energy resolution of about 0.1 eV.The data was collected in total electron yield (TEY) mode with thesampling depth being less than 10 nm (Abbate, et al., Surf. InterfaceAnal. 1992, 18, 65-69). The samples are loaded onto a stage beforetransferring them to the vacuum chamber via load lock. The pressureinside the sample chamber was maintained around 2×10⁻⁹ torr. Thevertical slit width used for these low energy XAS measurements was 100μm for Ti and O; and 50 μm for Ni L edge to enhance the resolution. Thespectra reported is obtained after averaging the data from multiplescans. TiO₂ and NiO powders were used as reference for calibrationpurposes. XANES data is normalized and analyzed using Athena software.

Gracing incidence XRD scans were performed on TiO₂:Ni (15 mol %) filmdeposited on Si(100) substrate. The crystal structure was observed to bedestroyed probably because of the formation of new NiTiO₃ phase andnon-uniformity of the film.

The Results of the Experiments Will Now be Discussed

TiO₂ films were spin coated onto Si substrates from an aged solutionproduced via sol-gel chemistry. The thickness was controlled based onspin speeds and number of coatings with the minimum thickness of a pureTiO₂ film measured to be 40±5 nm (Table 2). Film thickness with Nidopants was expected to be similar based on the processing but could notbe quantified due to the formation of a non-conformal film as observedvia AFM (FIG. 20).

TABLE 2 Thickness of the TiO₂ thin film measured using Filmetrics ToolDilution Spin Thin Film Ratio Speed Thickness TiO₂ — 3000 rpm 89.53 nmTiO₂ 1:1 3000 rpm 51.09 nm TiO₂ 1:2 3000 rpm 37.74 nm

To verify the Ni incorporation, UV-Vis absorption spectra was collectedfor pure TiO₂ and TiO₂:Ni sol deposited on glass substrates (Thickness˜500 nm) before annealing. Characteristic Ni²⁺ absorption peaks wereidentified, as shown in FIG. 21A (Donegan, et al., J. Lumin. 1986, 35,57-63), and assigned to their respective electronic transitions betweenthe t_(2g) and e_(g) levels according to theory for d⁸ electronicconfiguration in octahedral symmetry (Tanabe, et al., J. Phys. Soc. Jpn.1954, 9, 753-766; Tanabe, et al., J. Phys. Soc. Jpn. 1954, 9, 766-779;Tanabe, et al., J. Phys. Soc. Jpn. 1956, 11, 864-877; Sugano, et al., J.Phys. Soc. Jpn. 1958, 13, 880-899) (FIG. 22, and inset of FIG. 21A). Anextracted 10 Dq value of 1.10 eV was obtained from the UV-Vis spectrum(Table 3), which is less than pure TiO₂ (10 Dq=1.8 eV) (Finkelstein, etal., X-Ray Spectrom. 2002, 31, 414-418) and is attributed to theamorphous nature of the film. FIG. 21B shows the GI-XRD pattern of thepure TiO₂ film annealed at 450° C. with all peaks indexed to anataseTiO₂ (JCPDS#12-1272). In TiO₂:Ni film (FIG. 23), the crystal structurewas observed to be amorphous, which was attributed to the inability tomeasure the non-conformal film, instead probing the native SiO₂ layer.Therefore, in order to probe the local crystal structure and thedistribution of Ni dopants in the TiO₂:Ni film after annealing,aberration-corrected HRTEM imaging was performed in conjunction withEELS chemical mapping. Bright field HRTEM images (FIG. 24A) indicate thepresence of two phases which were identified as TiO₂ and an oxide phaseformed by Ni and Ti (yellow outline). Lattice spacings of 3.54 Å and1.91 Å were identified for (101) and (020) planes of anatase TiO₂ (FIG.25) (Ding, et al., J. Mater. Chem. 2011, 21, 1677-1680). Furthermore,the slight expansion observed in the lattice of TiO₂ is attributed tothe larger Ni ion incorporated as a substitutional dopant (Karthik, etal., J. Alloys Compd. 2011, 509, 5173-5176). The yellow outlined regionis the Ni dense phase in the TiO₂ film. The lattice fringe spacing inthis region was extracted as 2.12 Å, which matches to the (002) plane ofNiTiO₃ (Posnjak, et al., Z. Kristallogr.-Cryst. Mater. 1934, 88,271-280). Although, the formation of NiTiO₃ phase is not favored attemperatures below 600° C. Karimipour, et al., Phys. Scr. 2011, 84,035702; Bellam, et al., RSC Adv. 2015, 5, 10551-10559; Qu, et al., J.Mater. Chem. 2012, 22, 16471-16476), the lower solubility of the dopant(Ni) and higher annealing temperatures (450° C.) resulted in theagglomeration and growth of the NiTiO₃ clusters in the TiO₂ network(Karimipour, et al., Phys. Scr. 2011, 84, 035702). The EELScompositional mapping (FIG. 24B, FIG. 24C, and FIG. 24D) show thechemical maps for Ni, Ti, and O elements in the highlighted region (Nidense) of FIG. 24A. The elemental maps also indicate the formation of amixed phase of Ni, Ti, and O and sparsely distributed Ni in TiO₂ film.Furthermore, the electronic characterizations such as XPS and XAS(discussed below) suggest the presence of NiTiO₃ clusters. In eithercase of TiO₂:Ni or NiTiO₃, the local symmetry of Ni remains unchanged,i.e., Ni dopant ions are bonded to six O ions in an octahedral (O_(h))symmetry (FIGS. 26 and 27).

TABLE 3 Energy of Electronic transitions in Ni²⁺ Transition Expt Calc³A₂-¹E 798 nm 776 nm ³A₂-³T₁(³F) 647 nm 680 nm

In order to investigate the Ni 3d-O 2p hybridization in the TiO₂:Nifilm, surface sensitive XPS and XAS were performed. Due to the nature ofthese methods, the resulting response is a statistical representation ofthe film and not an individual atom, describing the overall chemical andphysical properties of the film. The survey scans (FIG. 28A), Ti 2p, O1s, and Ni 2p detailed spectra for the pure TiO₂ and TiO₂:Ni are shownin FIG. 28. The O 1s XPS spectra (FIG. 28B) is comprised of a stronglattice oxygen peak (O₁, Ti—O-TM; Bellam, et al., RSC Adv. 2015, 5,10551-10559) along with the shoulder peak (O₂). The positive shift inO₁, is attributed to the presence of a higher electronegative dopant(Ni²⁺) in the TiO₂ lattice (Roy, et al., J. Phys. Chem. C 2014, 118,29499-29506), suggesting the formation of a —Ni—O—Ti— bond. The areaunder the shoulder peak (O₂), ascribed to the oxygen bonded to undercoordinated cations, is observed to increase with Ni doping. When Ni²⁺substitutes Ti⁴⁺ ions, the charge in the lattice is compensated by thecreation of oxygen vacancies (Bharti, et al., Sci. Rep. 2016, 6, 32355;Bellam, et al., RSC Adv. 2015, 5, 10551-10559) resulting in undercoordinated cations as shown in the EELS maps (FIGS. 24B to 24D).Additionally, this peak can be attributed to surface hydroxyl (—OH)groups bonded to the hydrophilic TiO₂:Ni surface (Yao, et al., ACS Appl.Mater. Interfaces 2010, 2, 2617-2622). Next, Ti 2p detailed spectra(FIG. 28C) is performed and the characteristic Ti⁴⁺ spin-orbitalsplitting, with energies at 458.8 eV (2p_(3/2)) and 464.4 eV (2p_(1/2)),was observed for both the pure and doped films (Sanjinés, et al., J.Appl. Phys. 1994, 75, 2945-2951). The broadening of the Ti 2p_(3/2) mainpeak in TiO₂:Ni films is due to the Ti³⁺ shoulder, which is againattributed to the charge compensation TiO₂:Ni lattice (Bharti, et al.,Sci. Rep. 2016, 6, 32355). Ni 2p detailed scans (FIG. 28D) indicate thepresence of Ni²⁺ (Ding, et al., RSC Adv. 2015, 5, 95478-95487). Thespin-orbital splitting energy was measured as 17.48 eV, lower than thatof NiO (18.4 eV) (Yin, et al., Solid State Commun. 2005, 135, 430-433),suggesting that the NiO clusters are not present. Furthermore, the peakposition of the Ni 2p_(3/2) main peak at 855.04 eV corresponds to eitherNi²⁺ doped in TiO₂ lattice or to NiTiO₃ (Yu, et al., Sci. Rep. 2015, 5,9561; Qu, et al., J. Mater. Chem. 2012, 22, 16471-16476; Uddin, et al.,Phys. Chem. Chem. Phys. 2017, 19, 19279-19288). The two satellite peaks(6 eV, 9.5 eV) in Ni 2p spectra are attributed to the screening effectsof various core-hole and ligand hole states by the 3d and 4s bands(Grosvenor, et al., Surf. Sci. 2006, 600, 1771-1779). Deconvolutedspectra are shown in FIG. 29 and tabulated in Table 4.

TABLE 4 Binding Energies and Area under the Peaks of Ti and O in pureTiO₂ and TiO₂:Ni (15 mol %) films Ti⁴⁺ 2p_(3/2) Ti⁴⁺ 2p_(3/2) Ti⁴⁺2p_(1/2) Ti⁴⁺ 2p_(1/2) Thin film Pos Area Pos Area O₁ Pos O₁ Area O₂ PosO₂ Area TiO₂ 458.47 eV 68.48% 464.17 eV 31.52% 529.77 eV 77.40% 531.17eV 22.60% TiO₂:Ni 458.58 eV 49.41% 464.49 eV 25.40% 530.03 eV 63.18%531.56 eV 36.82% (15 mol %)

In order to understand the effect of the external dipole moment on theTM hybridization in a solid host, the as-prepared TiO₂:Ni films werefunctionalized with p-substituted benzoic acid ligands via carboxylicacid chemistry (Goh, et al., Appl. Phys. 2007, 101, 114503). Thesecarboxylic acids bridge to the surface of TiO₂ in a bidentate fashionresulting in a dipole moment normal to the surface, changing theelectron affinity of the organic (ligand)-inorganic (film) interface(Jankovic, et al., J. Phys. Chem. C 2009, 113, 12645-12652). The C—Hstretches were identified with FTIR (FIG. 30) for both the ligand-bondedsamples around 2800-3000 cm⁻¹ indicating that the organic group wasbound to the surface (Sundaraganesan, et al., Spectrochim. Acta, Part A2007, 67, 287-297). Furthermore, the NO₂ symmetric stretching mode wasalso observed at 1370 cm⁻¹ for NO₂-BZA functionalized sample (Osawa, etal., J. Phys. Chem. 1991, 95, 9914-9919).

After surface modification, the Ti ²p spectra was expected to remainunchanged as the spectra is dominated by the bulk of the crystal,whereas the adsorbate dipole moment acts near the surface (Rajh, et al.,J. Phys. Chem. B 2002, 106, 10543-10552). The 2p_(3/2) and 2p_(1/2) mainpeaks (FIG. 31) of Ni in TiO₂:Ni shifted less than half an eV,suggesting no change in the oxidation state after ligand bonding. Theasymmetric 6 eV satellite peak is due to the screening of the two-holestate, core hole and 3d hole, by the wide 4s band (Grosvenor, et al.,Surf. Sci. 2006, 600, 1771-1779). Furthermore, the interatomic wavefunction mixing of the Ni 3d states and the ligand p states influencethe screening of these multiplet effects. In the present case of surfacemodified TiO₂:Ni films, the aromatic surface ligands are expected todelocalize the hole wave functions, i.e, conduction band 3d orbitals ofthe inorganic layer (Teunis, et al., Nanoscale 2017, 9, 14127-14138).Despite the slight changes in the 6 eV satellite peak positions of theligand-bonded TiO₂:Ni films, it is hard to decipher the effect ofelectronegativity of the ligand on the core states of Ni 2p, due tovarious factors such as instrument resolution, interatomic wave functionmixing (Ni 3d-O 2p) (Jana, et al., Chem. Mater. 2016, 28, 1107-1120),and atomic multiplet coupling. Ni 2p deconvoluted spectra are shown inFIG. 32 and tabulated in Table 5.

TABLE 5 Binding Energies of Ni²⁺ XPS Spectra in TiO₂:Ni (15 mol %) filmsThin Film 2p_(3/2) 2p_(s,3/2) 2p_(1/2) TiO₂:Ni 855.05 eV 861.33 eV872.53 eV NH₂-BZA 855.25 eV 861.55 eV 872.73 eV NO₂-BZA 855.62 eV 861.82eV 873.06 eV

The XPS spectra of Ti 2p and O 1s in TiO₂:Ni films indicate that theformation of the new oxide phase NiTiO₃ has not affected the structureof TiO₂, suggesting that the surface of TiO₂:Ni film has octahedrallycoordinated Ni²⁺ and Ti⁴⁺ ions. However, there is no strong evidencefrom XPS that suggests the perturbation of site symmetry or core levelenergy of these ions upon ligand bonding. Therefore, in order tounderstand the subtle changes in the electronic and geometric structureof the inorganic film at the hybrid interface, XANES was employed. The OK edge spectra (FIG. 10A) demonstrates the electronic transitions from O1s to the derived states of O 2p. The t_(2g) and e_(g) splitting of theTi 3d states hybridized with O 2p states is evident in the low energyregion of TiO₂:Ni film, whereas the higher energy spectral featuresindicated the formation of Ni—O bond. The broadening of these spectralfeatures also indicates the absence of long-range order of the locallattice (de Groot, et al., Phys. Rev. B 1989, 40, 5715), as observed inXRD. FIG. 10B shows the TEY spectrum of Ti L_(III) (2p_(3/2)-3d, 460.8eV) and L_(II) (2p_(1/2)-3d, 468.4 eV) edges of pure and TiO₂:Ni films.The doublet of the e_(g) band in TiO₂ L_(III) edge is attributed to thenon-degenerate d_(z) ₂ and d_(x) ₂ _(-y) ₂ , which is a signature ofanatase crystal structure (Henderson, et al., Phys. Chem. Miner. 2002,29, 32-42). This splitting dampens upon Ni doping due to the oxygenvacancies altering the Ti—O bonding environment, causing non-cubicstructural distortion in TiO₂:Ni film (Chen, et al., Phys. Chem. Chem.Phys. 2015, 17, 22064-22071; Radtke, et al., Phys. Rev. B 2006, 74,155117). Moreover, the onset of the absorption edge for TiO₂:Ni films isshifted to lower binding energies due to the presence of Ti³⁺ ions asobserved in Ti 2p XPS of TiO₂:Ni (15 mol %) films (Stoyanov, et al., Am.Mineral. 2007, 92, 577-586). After surface modification, the O K edge(FIG. 10C) and Ti L_(III/II) edge (FIG. 10D) spectra for TiO₂:Ni filmsindicated a change in the crystal field splitting energy (10 Dq). TheXANES spectra of Ti and O in the surface modified TiO₂ films did notshow any difference with surface functionalization due to the weakpenetration of these ligand fields. The energy splitting between thet_(2g) and e_(g) hybridization peaks in the O K edge of TiO₂:Ni films isproportional to the adsorbate dipole moment of the ligand (FIG. 33A)(Tsuchida, Bull. Chem. Soc. Jpn. 1938, 13, 388-400). In addition to thespectral differences in the lower energy region, the higher energyspectral features of O K edge, which correspond to the O 2phybridization with the Ti/Ni valence levels also indicate significantdifferences with the ligand. The onset of the second peak in the O 1s-O2p (4sp) region suffered a shift of about 0.1 eV from the referenceTiO₂:Ni (15 mol %), suggesting a change in the hybridization of the O 2porbitals with the ligand. Moreover, the broadening of these peaks forthe NO₂-BZA bonded films, indicate the covalent nature of themetal-oxygen bond accompanied with slight geometric distortion at thehybrid interface. In the Ti L edge spectra, while the Ti L_(III/II) edgepeak positions are indicative of the crystal field splitting, de Grootand co-workers showed that the FWHM of the e_(g) peak in the Ti L_(II)edge is also proportional to the 10 Dq (de Groot, et al., Phys. Rev. B1990, 41, 928-937). The deconvoluted Ti L_(II) edge spectra (FIG. 33B,FIG. 33C, and FIG. 33D) show a systematic increase in the FWHM of thee_(g) peak with an increasing ligand dipole moment. The ΔE in O K edgepeak positions and FWHM of the e_(g) peak in Ti L_(II) edge aretabulated in Table 6.

TABLE 6 ΔE of O K edge and FWHM of e_(g) peak in Ti L_(II) edge forsurface modified TiO₂:Ni (15 mol %) films, indicating the crystal fieldsplitting shifts as a function of the surface dipole moment. FWHM e_(g)Thin Film ΔE (eV) (L_(II) Edge) TiO₂:Ni 2.36 2.11 NH₂-BZA 2.28 2.01NO₂-BZA 2.44 2.35

To acquire the fingerprint analysis on the 3d electronic states of Niwhich are influenced by surface ligands, XANES spectra for Ni K andL_(III/II) edges in the TiO₂:Ni film were collected. The K edge XANESspectra was compared with standard NiO reference powder (FIG. 34) toidentify the oxidation state of Ni as 2+ in TiO₂:Ni film. The NiL_(III/II) edge spectra (FIG. 35A) for ligand bonded TiO₂:Ni films aresplit into L_(III) and L_(II) edges due to 2p spin-orbit coupling. Thetwo peaks in L_(II) edge correspond to the 3d states (t_(2g), e_(g)) ofNi²⁺ ion bonded to O²⁻ in O_(h) symmetry (Radtke, et al., Phys. Rev. B2006, 74, 155117; Matsuo, et al., Anal. Sci. 2001, 17, 149-153). A clearshift is observed in the L_(II) edge t_(2g)/e_(g) peak intensities as afunction of the adsorbate dipole moment. This difference is attributedto the change in the Ni e_(g)-O p hybridization with the ligands (deGroot, et al., Phys. Rev. B 1990, 41, 928-937; Guo, et al., Phys. Chem.Chem. Phys. 2016, 18, 3250-3259). The well-resolved multiplet structureon the L_(III) and L_(II) edges of NH₂-BZA bonded films compared to thatof NO₂-BZA films suggest the ionic nature of the Ni-bond (Wang, et al.,J. Am. Chem. Soc. 2000, 122, 10544-10552), which is attributed to theweaker metal-oxygen orbital overlap.

The L_(III/II) edge XANES spectra of Ni²⁺ in TiO₂:Ni was simulated usinga ligand field dependent simulation software, CTM4XAS (Stavitski, etal., Micron 2010, 41, 687-694; de Groot, et al., Phys. Rev. B 1990, 42,5459). These structures have been modeled in past with distortedsymmetries (D_(3d)/D_(2d)) demonstrating no difference with the O_(h)symmetry. The key parameters involved in these calculations are ligandfield parameter (10 Dq), charge transfer energy (Δ), Hubbard core-holepotentials (U_(pp), U_(pd)), slater integrals (F_(pp), F_(pd)), and thehopping parameters ((T(t_(2g)), T(e_(g))). The values for theseparameters were obtained from previous calculations and are summarizedin Table 7. The value of 10 Dq used in these calculations was obtainedfrom the UV-Vis absorption spectra. Furthermore, it was observed thatchanging the value of 10 Dq does not simulate the experimentallyobserved variation in Ni L_(II) edge (t_(2g)/e_(g)) branching ratio.Therefore, to understand the effect of electronegative/electropositiveligand on the hybridization of the Ni²⁺ 3d states, all the parametersexcept the hopping parameters were based on literature. The NiL_(III/II) edge spectra of NH₂-BZA bonded TiO₂:Ni film (FIG. 35B) wasmodeled using standard hopping parameters (T(e_(g))=1.8, T(t_(2g))=1) inO_(h) environment (de Groot, Coord. Chem. Rev. 2005, 249, 31-63).Simulations show that shift in the onset of the absorption edge inNH₂-BZA bonded TiO₂:Ni films to higher energies is due to the spinexchange interactions. The Coulomb (U) and charge transfer energies (Δ)increase for electron donating ligands (NH₂-BZA) bonded TiO₂:Ni. This isattributed to an electron addition to the high-spin Ni²⁺ 3d⁸ staterequiring additional spin exchange stabilization energy, pushing theleading absorption edge to higher energies (Qiao, et al., Curr. Appl.Phys. 2013, 13, 544-548). On the other hand, for NO₂-BZA bonded films,the simulated spectra matched the experimental results by consideringstrong ligand character in the ground state 3d⁸ L orbitals. This is doneby setting the values of the hopping parameters of t_(2g) and e_(g)equal to 1.8, indicating strong mixing between the t_(2g), e_(g)orbitals (Stavitski, et al., Micron 2010, 41, 687-694; de Groot, et al.,Phys. Rev. B 1990, 42, 5459; de Groot, Coord. Chem. Rev. 2005, 249,31-63). Furthermore, these non-standard values of hopping parameterssuggest geometric distortion at the interface, i.e. the transformationof an octahedron (FIG. 35C, left) into a low symmetry structure such assquare planar, as shown in FIG. 35C, middle. From an experimentalstandpoint, the mixing of the t_(2g) and e_(g) energy levels in NO₂-BZAbonded films can be attributed to the strong covalent nature of the Ni—Obond. Additionally, the increased intensity of the e_(g) peak in NO₂-BZAbonded films corresponds to the reduced electron density in the Ni 3dstates. Both of these effects were observed in the O K edge spectra ofthe NO₂-BZA bonded films (FIG. 10C). However, as the model suggests, thetransformation of an octahedron to square planar structure involvesdistinct structural changes accompanied by oxygen vacancies (Wu, et al.,J. Phys. Chem. C 2012, 116, 7219-7226), which was not observed with theother characterization techniques. Therefore, the observed branchingratio in NO₂-BZA bonded TiO₂:Ni film can be interpreted as a slightbreaking of the O_(h) symmetry due to the elongation along the axial(e_(g) orbital) direction of the Ni atom as shown in FIG. 35C, right.

TABLE 7 Values of Variables in the CTM4XAS Calculation Parameters SymbolValue Slater Integrals Fpd, Fdd, Gpd 0.8, 0.8, 0.8 Symmetry O_(h) O_(h)Crystal Field Splitting 10 Dq  1.1 eV Spin Orbital Coupling SO couplingreduction 1 Charge Transfer Δ    3 eV Hubbard Potentials Udd, Upd 6, 8eV Hopping Parameters T(e_(g)), T(t_(2g)) 1.8, 1; 1.8, 1.8

The interplay between the surface dipole, electronic states, p-dhybridization, and the crystal field splitting energy of Ni²⁺ 3dorbitals is illustrated in FIG. 36. The electronic structure of TiO₂:Ni(FIG. 36A) shows the formation of interband gap Ni 3d states, whichresulted in lowering of the Fermi energy (Matsumoto, et al., J. Ceram.Soc. Jpn. 2010, 118, 993-996). The energy of these interband gap Ni 3dstates is manipulated by the ligand due to hybridization (FIG. 36B) withthe neighboring O atoms resulting in a shift in orbitals energyproportional to the crystal field splitting. This ability to tune thelocal crystal field allows for control of the optical states via themodified hybridization of TM e_(g) orbitals as shown in FIG. 36C.Unfortunately, the shift was not detectable in the UV-Vis measurementsdue to the weak absorption of the thin, non-conformal films. However,the surface sensitive XAS measurements identified the change in theelectron density in the valence 3d states of Ni²⁺ with the ligand.Effectively, the surface dipole manipulates the energy of Ni 3d morethan bulk Ti 3d orbitals, impacting the electron density in thoseorbitals. In particular, the e_(g) orbitals, which are pointed towardsX, Y, and Z axes (de Groot, et al., Phys. Rev. B 1990, 41, 928-937)overlap with the p orbitals of the neighboring O atoms, wherein thestrength of the hybridization is determined by the nature of the dipole.As seen with the XAS spectra, the change in the hybridization of theNi—O bond in TiO₂:Ni film is quantified by the shift in the Ni L_(II)edge branching ratio (t_(2g)/e_(g)), FWHM of the e_(g) peak in the TiL_(II) edge, and the ΔE of the O K edge. This congruency among thevalence level spectra of all the elements reinforces the fact that theelectronic structure of TM dopant in a solid host is a function of thedipole moment of the surface ligand.

In summary, TiO₂:Ni films were spin-coated on Si (100) substrates usingsol-gel chemistry. Initial structural and optical characterizationresults confirmed the crystal structure and crystal field splittingenergy. The crystalline nature of the TiO₂:Ni (15 mol %) film waslocally determined using HRTEM to identify the presence of TiO₂:Ni andNiTiO₃ phases. Furthermore, spatial mapping of these films using EELSconfirmed those two phases. Surface functionalization of these inorganicfilms was performed with weak benzoic ligands via carboxylic acidchemistry to apply a weak external field in the form of a dipole moment.The surface dipoles were observed to show no effect on the pure TiO₂films owing to the bulk-like characteristics of the elements present inthe film. The influence of the surface dipole on core and valenceelectronic states of the TM dopant in TiO₂:Ni²⁺ was systematicallyinvestigated by surface sensitive characterization techniques such asXPS and XAS. The results from these characterization methods point tothe change in the ligand character of the Ni 3d orbitals. It isimplicitly proven that the overlap between the Ni 3d orbitals and O 2porbitals is a function of the dipole strength of the surface ligand.This ability to control the hybridization of TM ion in a solid host viaweak external fields can be utilized to engineer the optical andmagnetic responses in a device. Specifically, the adaptive opticalproperties of TM doped solids can be coupled with the steady rare earthemissions in inorganic phosphors to obtain dynamic luminescence andthereby, minimize the usage of multiple rare earth doped Red-Green-Blue(RGB) phosphors. Furthermore, these hybrid luminescent materials due totheir tunable properties will have potential applications in flexibleelectronics, biosensors, and solar cells.

Example 5: Effect of Moisture on Dopant Segregation in Solid Hosts

TiO₂ is one of the most extensively studied compounds for a myriad ofapplications such as photocatalysis (Peng, et al., Advanced PowderTechnology 2012, 23, (1), 8-12; Schneider, et al., Chemical reviews2014, 114, (19), 9919-9986), solar cells (Park, et al., The Journal ofPhysical Chemistry B 2000, 104, (38), 8989-8994; Frank, et al.,Coordination Chemistry Reviews 2004, 248, (13-14), 1165-1179), andelectrochemistry (Ding, et al., RSC Advances 2015, 5, (116),95478-95487; Liu, et al., International Journal of Hydrogen Energy 2015,40, (5), 2107-2114), owing to its non-toxicity (Hafizah, et al.,International Journal of Photoenergy 2009, 2009), oxidation power (Su,et al., Applied Catalysis B: Environmental 2008, 77, (3-4), 264-271),chemical stability (Inturi, et al., The Journal of Physical Chemistry C2013, 118, (1), 231-242), and resistance to corrosion (Bharti, et al.,Scientific reports 2016, 6, 32355; Macak, et al., small 2007, 3, (2),300-304). The relatively high band gap of TiO₂, about 3.2 eV (Scanlon,et al., Nature materials 2013, 12, (9), 798), unfortunately, limits itsphotocatalytic/photovoltaic effect to the UV region, which constitutesonly 5% of the solar spectrum (Qin, et al., Chemical communications2010, 46, (13), 2304-2306). Additionally, the fast recombination rate ofphoto-generated electrons and holes⁴ also reduce the photonic efficiencyof TiO₂. Therefore, in order to extend the optical absorption of TiO₂from the UV to the visible region (Pelaez, et al., Catalysis Today 2009,144, (1), 19-25; Asahi, et al., Science 2001, 293, (5528), 269-271; Ren,et al., Applied Catalysis B: Environmental 2007, 69, (3-4), 138-144),while simultaneously improving the photocatalytic performance, variousmethods such as chemical doping (Liu, et al., International Journal ofHydrogen Energy 2015, 40, (5), 2107-2114; Li, et al., Physical ChemistryChemical Physics 2013, 15, (46), 20037-20045), dye sensitization (Pan,et al., Journal of the American Chemical Society 2011, 133, (26),10000-10002; Ni, et al., Renewable and Sustainable Energy Reviews 2007,11, (3), 401-425), and narrow gap semiconductor coupling (Liu, et al.,international journal of hydrogen energy 2011, 36, (1), 167-174; Chen,et al., Chemical reviews 2007, 107, (7), 2891-2959) have been employedto modify the electronic structure. Among all of these, doping withmetals, in particular transition metal (TM) elements, has received greatattention due to the effective narrowing of the band gap. The band gapis narrowed in these homogeneous TM doped TiO₂ materials by theformation of interband gap states that can also act as trapping sitesfor preventing the carrier combination of photo-generated chargecarriers (Wang, et al., Nanoscale 2012, 4, (21), 6682-6691).

While there have been many studies that have focused on optimizing thedoping concentration to yield select photophysical responses (Rodrigues,et al., Eclética Química 2011, 36, (1), 18-36), the positioning of thedopant (substitution vs. interstitial, surface vs. bulk) in the hostlattice still remains a synthetic challenge. The spatial distribution ofthe dopants, which is often not emphasized in literature (Xie, et al.,Journal of Power Sources 2013, 224, 168-173; Kumar, et al., Thin SolidFilms 2016, 619, 144-147), is an important parameter that determines thephysical, chemical, optoelectronic, and magnetic properties of the TMdoped TiO₂. There have been few studies exploring this argument; forexample, the visible light photocatalytic activity of N doped TiO₂(Peng, et al., Journal of Solid State Chemistry 2008, 181, (1), 130-136)and charge carrier concentration in CdSe nanocrystals (Sahu, et al.,Nano letters 2012, 12, (5), 2587-2594) are two instances where theproperties of the material are sensitive to the dopant position in thehost lattice. Additionally, in all of the above mentioned literaturereports, the homogeneous distribution of dopants in the host TiO₂lattice is validated based on bulk characterization methods such as XRD,UV-Vis, and EDX that are insensitive to the local dopant environment inthe host lattice (Navas, et al., Physical Chemistry Chemical Physics2014, 16, (8), 3835-3845; Karlsson, et al., Toxicology letters 2009,188, (2), 112-118; Tripathi, et al., Adv Mater Lett 2015, 6, 20).Furthermore, the resultant optoelectronic properties in those dopedsolids are elucidated based on the electronic band structure for thecase of where TiO₂ is homogeneously doped. However, these discussions donot address the issues of dopant segregation in the host lattice and itsimpact on the derived properties. At high doping concentrations, beyondthe solubility limit of the dopant in the host lattice, the formation ofmetallic/metal oxide clusters is expected (Mesilov, et al., The Journalof Physical Chemistry C 2017, 121, (43), 24235-24244). Additionally, thelocation of these dopant/dopant oxide clusters, including their surfaceadsorbates (hydroxyls), determine the catalytic properties of thesystem.

In the present work, Ni²⁺ doped TiO₂ nanoparticles (NPs), with dopingconcentration upto 15 mol %, were synthesized via sol-gel chemistry. Thehigh dopant concentration provides more catalytically active sites forsolar-driven applications. Secondary processing conditions of the sol,such as drying and annealing, were observed to influence the segregationof NiO clusters in TiO₂, altering the resulting optoelectronicproperties. The bulk crystal structures of TiO₂:Ni (15 mol %) NPs wereidentified from XRD and the evolving local environment around the dopantwas probed via HRTEM images coupled with elemental chemical mapping. Theaging of the dried TiO₂:Ni (15 mol %) powders was observed to be amoisture sensitive phenomenon, which was studied using time resolvedUV-Vis and FTIR spectroscopy measurements. TGA-DSC studies wereperformed on the aged and non-aged TiO₂:Ni (15 mol %) dried powders toelucidate the dopant segregation mechanism in the aged powders.Furthermore, in order to control the NiO cluster formation in the hostTiO₂ matrix, rapid annealing was performed on TiO₂:Ni (15 mol %) driedpowders. The doped system was frozen into metastable excited states byquenching, and as a result, the dopants are locked in the host latticesites. Similar trend of dopant segregation was observed with other firstrow TM (Co²⁺) doped TiO₂ powders upon slow annealing. This work extendsour previous X-ray local structure studies on 40 nm thick TiO₂:Ni films(Darapaneni, et al., The Journal of Physical Chemistry C 2018) intoanother diminished dimension of nanoparticles (˜20 nm dia), allowing forbetter understanding of the fundamental processes such as dopantincorporation and segregation in host matrices, to establish thestructure-processing-property relationship. This ability to incorporatehigher dopant concentrations while controlling the dopant position inhost lattice can be utilized to engineer the dopant effects in the host,which will assist in the development of solar cells, photocatalyticdevices, etc. with improved optoelectronic properties andphotoconversion efficiencies.

Materials and Methods

Titanium (IV) Isopropoxide (TTIP, Acros Organics, >98%), nickel (II)chloride hexahydrate (NiCl₂.6H₂O, BTC, >99%), cobalt (II) chloridetetrahydrate (CoCl₂.4H₂O, Sigma Aldrich), hydrochloric acid (HCl,36-38.5% purity, ACS grade), reagent alcohol (<0.075% VWR Analytical)were obtained commercially. All the materials were used without furtherpurification.

Ni Nanoparticles (NPs): TiO₂:Ni (15 mol %) NPs were synthesized byemploying the facile sol-gel chemistry (Yu, et al., Scientific reports2015, 5, 9561). The sol was prepared by dissolving 207 mg of NiCl₂.6H₂Oin 5 mL of ethanol and then adding 1.5 mL of TTIP dropwise undervigorous stirring. Homogeneous TiO₂:Ni sol was obtained after 3-4 h ofcontinuous stirring. HCl (125 μL) was used as a catalyst in this processto control the rapid hydrolysis of TTIP precursor. The dopingconcentration of Ni precursor to TTIP was varied from 0 to 15 mol %. Theprepared sol was aged for 24 h before drying it in air for 10 h. Thesedried powders were further aged in air for 0-48 h to systematicallystudy the effect of moisture on the dried TiO₂:Ni (15 mol %) powders.Similarly, TiO₂:Co (15 mol %) NPs were synthesized with CoCl₂.4H₂O asprecursor. Pure TiO₂ NPs were prepared in the same method without theaddition of Ni/Co precursor. All the dried powders were annealed in airat 450° C. for 2 h to form crystalline powders.

Aberration corrected STEM images were obtained using the 100 kVNion-UltraSTEM 100 (U100) electron microscope equipped with thirdgeneration C₃/C₅ aberration corrector, Gatan Enfina electron energy lossspectrometer (EELS), and a cold FEG source. The crystal structure wasidentified by performing Powder X-ray Diffraction (XRD) usingPANalytical X-ray diffractometer operating at 45 kV and 40 mA. The θ-2θradial scan was performed over the range 5-70° with a step size of 0.03°and dwell time of 60 s, using Cu K_(α1) (λ=1.54 Å) as radiation source.Thermogravimetric analysis (TGA) was performed with a TA SDT Q600DSC-TGA under air flow to understand the crystallization process ofdried TiO₂:Ni powders. The temperature was programmed from 25 to 450° C.at 4° C./min, held 120 min, and then cooled to room temperature at thesame rate of 4° C./min.

The absorption spectra of TiO₂:Ni NPs was recorded using a Perkin-ElmerLambda 900 UV/Vis/NIR spectrometer equipped with an integrating sphereand a center-mounted sample holder. The absorption scans ranging from300 to 1300 nm with a scan rate of 1 nm/s were obtained on the NPs driedon the glass substrates. The change in monochromators was set to occurat 900 nm. Fourier-transform infrared (FTIR) spectroscopy was performedon these TiO₂:Ni NPs using DRIFTS mode of measurement in a ThermoScientific Nicolet 380 FTIR with a DTGS detector. The data was collectedin Kubelka-Munk (f(R)) mode, with air as background, 30 min of N₂ purge,and resolution being 4 cm⁻¹ in the region going from 4000 to 1000 cm⁻¹.

The Results of the Experiments Will Now be Discussed

TiO₂:Ni (15 mol %) NPs were prepared by drying the aged sol followed byannealing at 450° C. for 2 h in air. These NPs are spherical in shapewith 20 nm average diameter. The XRD patterns of the annealed TiO₂:Ni(15 mol %) NPs are shown in FIG. 14A. A clear difference with respect tothe formation of NiO was observed with the aging time of the dried NPsprior to annealing. TiO₂:Ni (15 mol %) NPs that are annealed directlyafter drying exhibited a doped anatase phase while the aged powdersafter drying resulted in the formation of segregated NiO phase. However,the detailed XRD scan of the non-aged and annealed TiO₂:Ni (15 mol %)NPs (inset of FIG. 14A), shows the presence of small NiO peak. It can beinferred that the formation of NiO clusters in TiO₂ is inevitable withhigh doping concentrations, i.e., 15 mol %, which exceeds the criticaldoping concentration of 10 mol % (Vakhitov, et al., Journal of Physics:Conference Series, 2014; IOP Publishing: 2014; p 012048; Chou, et al.,Applied Energy 2014, 118, 12-21; Bilecka, et al., The Journal ofPhysical Chemistry C 2011, 115, (5), 1484-1495). In addition, theintensities of the XRD peaks for both the NP samples suggests that thesize of these segregated NiO clusters is a function of the aging time.Therefore, to differentiate the effect of these two parameters,concentration and aging time, on NiO segregation, control experimentswere performed on 5 and 10 mol % Ni doped TiO₂ NPs. Detailed XRD scanswere taken on the TiO₂:Ni (5 and 10 mol %) NPs which were aged andnon-aged prior to annealing. The non-aged powders show no traces of NiOclusters whereas the aged 10 mol % powder exhibit small NiO peak. Withthese experiments, it is evident that aged TiO₂:Ni powders form NiOclusters upon annealing, due to the interaction of the dopant (Ni) inTiO₂ organic matrix with atmospheric moisture. Furthermore, the UV-Visabsorption measurements were performed on these aged and non-agedannealed NPs to corroborate the XRD observations and identify the Nicoordination in the TiO₂ matrix. The spin forbidden transitions such as³A_(2g)-¹T_(1g) (G) and ³A_(2g)-¹E (D) are intense in NiO NPs due to thespin-orbit coupling and antiferromagnetic order in bulk NiO.⁴¹Similarly, the TiO₂:Ni NPs with NiO segregation showed distinct peaksattributed to spin forbidden ³A_(2g)-¹E (D) transitions while thenon-NiO segregated TiO₂:Ni NPs exhibited broad peaks of spin-allowedtransitions (Brik, et al., ECS Journal of Solid State Science andTechnology 2016, 5, (1), R3067-R3077).

To better understand the local structural differences with respect tothe NiO formation in the aged versus non-aged TiO₂:Ni (15 mol %)powders, spatially resolved bright-field and dark-field HRTEM imageswere taken. The lattice fringe spacings extracted for the aged andannealed TiO₂:Ni (15 mol %) NPs indicate two regions: TiO₂ and NiO. Thedopant is observed to segregate as large NiO cubic clusters withoutforming a doped TiO₂ phase. The corresponding EELS spectra in thoseregions also indicates the absence of Ni/Ti peaks in the TiO₂/NiO richregion. Similarly, the non-aged and annealed TiO₂:Ni (15 mol %) NPs alsoexhibit two regions in the bright-field images: TiO₂ and NiO as shown inFIG. 14B. The lattice planes for anatase TiO₂ were observed to be thedominant group in the sample with smaller clusters of NiO. Moreover, theEELS spectra on these non-aged and annealed NPs suggest the presence ofNi in both TiO₂ and NiO phases. These results further confirm that theformation of large NiO clusters in aged TiO₂:Ni (15 mol %) NPs is due tothe strong interaction of the atmospheric moisture with the dopant (Ni)in the TiO₂ organic matrix.

To investigate the effect of atmospheric moisture on air exposed TiO₂:Ni(15 mol %) dried powders, Fourier Transform Infrared Spectroscopy (FTIR)was performed. FIG. 15A shows the FTIR spectra of the aged and thenon-aged TiO₂:Ni (15 mol %) powders after drying. The O—H stretches forthe terminally bound and bridged hydroxyl groups were present in the3300-3700 cm⁻¹ region (Finnie, et al., Langmuir 2001, 17, (3), 816-820).The peak at higher wavenumbers, around 3540 cm⁻¹, corresponds to theterminally bound OH species, which often hydrogen bond with the adsorbedwater molecules. A sharp increase in this peak was observed for the agedTiO₂:Ni (15 mol %) powders, and is attributed to physisorbed watermolecules upon atmospheric moisture exposure (Anpo, et al.,Environmentally benign photocatalysts: applications of titaniumoxide-based materials. ed.; Springer Science & Business Media: 2010).For the pristine dried powders, distinct peaks corresponding to the C—Halkane stretches (2930 and 2970 cm^(−l)), C—O stretching bands (1130cm⁻¹) and CH₂ bending bands (1130 cm⁻¹ and 1380 cm⁻¹) were identified(Kratochwil, et al., Journal of electron spectroscopy and relatedphenomena 1993, 64, 609-617; Rasko, et al., Applied Catalysis A: General2004, 269, (1-2), 13-25; Dobson, et al., Spectrochimica Acta Part A:Molecular and Biomolecular Spectroscopy 1999, 55, (7-8), 1395-1405).However, all of these peaks diminished in intensity after aging for 48 hin air. While there are many ethoxide and ethanol groups adsorbed on thesurface of dried TiO₂, it has been shown that the ethoxy group isdestabilized in the presence of water and leaves the metal oxide surfaceas ethanol. Therefore, the dampening of the C—H peaks after aging in airis attributed to the desorption of the ethoxy group in the presence ofmoisture. Additionally, the Ti—O—C peak at 1040 cm⁻¹, which correspondsto the unhydrolyzed alkoxide species also disappeared for aged samples,suggesting a change in the organic framework upon aging (Pillai, et al.,The Journal of Physical Chemistry C 2007, 111, (4), 1605-1611; Zywitzki,et al., Journal of Materials Chemistry A 2017, 5, (22), 10957-10967).Similar trend of increase in the concentration of physisorbed water anddisappearance of alkane stretches/bending was observed for the airexposed pure TiO₂ powders. These FTIR results clearly demonstrate thatthe dried TiO₂:Ni (15 mol %) powders are undergoing hydroxylation uponatmospheric exposure, and as a result, the surface hydroxylconcentration is increased for the aged powders.

Complementing the FTIR studies, time-resolved UV-Vis absorptionmeasurements were performed on TiO₂:Ni (15 mol %) dried powders prior toannealing for systematic investigation of the changes induced in thebond structure of TiO₂:Ni matrix upon moisture exposure. The driedpowders were exposed to atmosphere and samples were collected every 3,6, 9, 12, 24, and 48 h. The UV-Vis absorption spectra of these powdersis shown in FIG. 15B. A color change from yellow to green was observedfor the dried TiO₂:Ni (15 mol %) powders upon aging in air from visualinspection. This is evident from the blue shift of the absorptionspectrum with increasing exposure time. The yellow to green color changeis indicative of a change in the crystal field splitting energy, 10 Dq.Moreover, previous studies on complex Ni oxides report that the brightyellow color is a result of distorted octahedral environment whereas thegreen color corresponds to octahedral symmetry around the Ni ion. Thischange in the value of 10 Dq and the corresponding change in the localsymmetry, point to the claim that the interaction of moisture with thedopant is causing a change in the bond structure around Ni in the hostTiO₂. The 10 Dq values calculated for the dried TiO₂:Ni powders exposedto moisture for different time intervals is tabulated in Table 8. Inaccordance with the blue shift of the absorption spectra, the increasein 10 Dq with increasing moisture exposure time, indicates the bondingof cation in TiO₂:Ni (15 mol %) powders to more electronegative hydroxylgroups (Huheey, The Journal of Physical Chemistry 1965, 69, (10),3284-3291) based on the spectrochemical series (Ryutaro, Bulletin of theChemical Society of Japan 1938, 13, (5), 388-400).

TABLE 8 Increase in the value of 10 Dq of dried TiO₂:Ni (15 mol %)powders with aging time Transition Time (h) ³A₂-³T₂ (F) 0 6 48 Expt. 10Dq 0.92 eV 0.97 eV 1.00 eV Calc. 10 Dq 0.94 eV 1.04 eV 1.13 eV

The UV-Vis and FTIR spectroscopic results suggest an increase in theconcentration of surface hydroxyl groups bonded to the cation uponaging. However, these results cannot predict the reaction mechanism forNiO segregation in air exposed TiO₂:Ni (15 mol %) powders. Therefore,TGA-DSC studies were performed on dried TiO₂:Ni (15 mol %) powders toelucidate the reaction kinetics of moisture-dopant interaction. Two setsof amorphous TiO₂:Ni (15 mol %) powders, one exposed to air for 48 hafter drying and other sealed in a vial after drying, were subjected tothermal treatment in TGA furnace at a constant heating rate of 4° C./minin the range of 25−450° C. Pure TiO₂ was employed as the control samplein these measurements. The heat flow for these samples was constantlyobserved using the DSC data. FIG. 37 shows the TGA-DSC data of driedTiO₂:Ni (15 mol %) and TiO₂ NPs. The first endothermic peak at 70° C. isascribed to the evaporation of the ethanol solvent (Chen, et al.,Applied Surface Science 2007, 253, (23), 9154-9158; Wu, et al., ChinaParticuology 2003, 1, (6), 262-265). The corresponding weight loss inthe temperature range of 25-125° C. is recorded as 29.9% and 3.4% forthe aged and non-aged TiO₂:Ni (15 mol %) powders respectively. Thehigher weight loss associated with the 48 h aged TiO₂:Ni (15 mol %)powders suggests the volatilization of ethanol from the hydrated TiO₂matrix. The increase in the weight loss in the lower temperature region(25-125° C.) with aging time was confirmed by performing similar heattreatment studies on the 3 h aged TiO₂:Ni (15 mol %) dried powders. Thisphenomenon of replacement of ethoxy with hydroxyl groups upon aging isconsistent with the previous results of UV-Vis and FTIR measurements. Atslightly higher temperatures of about 100° C., the evaporation ofphysisorbed water (Finnie, et al., Langmuir 2001, 17, (3), 816-820) wasobserved in the DSC plot for the aged TiO₂:Ni (15 mol %) powders.However, in pure TiO₂, the volatilization of organic solvent and wateris combined in a broad endothermic region extending from 50 to 220° C.With further increase in temperature, the non-aged TiO₂:Ni (15 mol %)powder exhibits a sharp exothermic peak at 230° C., similar to the pureTiO₂, indicating the combustion of the organic compounds (Muniz, et al.,Ceramics International 2011, 37, (3), 1017-1024). The crystallization ofthe amorphous powder into anatase TiO₂ is observed for all the samplesat 400° C.⁷ For the aged TiO₂:Ni (15 mol %) powders, a broad exothermicpeak is observed from 220-400° C. The first region in between 220-330°C. is attributed to the removal of the unhydrolyzed isopropoxide groups(Hafizah, et al., International Journal of Photoenergy 2009, 2009)whereas the second region from 330-400° C. corresponds to the formationof nickel oxide phase (Wu, et al., Materials Letters 2007, 61, (14-15),3174-3178; Li, et al., Materials Letters 2007, 61, (8-9), 1615-1618).This co-existence of the NiO and anatase TiO₂ phases in aged TiO₂:Ni (15mol %) NPs was observed earlier in the XRD patterns (FIG. 14).

On the basis of above observations, the dopant incorporation andsegregation mechanism in solid hosts is illustrated using the schematicshown in FIG. 38. The doping of Ni in TiO₂ follows a nucleation-dopingmechanism, as the Ti—O—Ni bonds are formed in the organic frameworkduring the polycondensation step of the sol-gel synthesis (Lynch, etal., The Journal of Physical Chemistry C 2015, 119, (13), 7443-7452;Lee, et al., Materials Science and Engineering: B 2006, 129, (1-3),109-115). After drying the sol in air for 10 h at 100° C., the amorphousTiO₂:Ni (15 mol %) powders are composed of an organic matrix, in whichthe unevaporated ethanol and water molecules bond dissociatively to thecations (Ti, Ni). Upon aging these powders in air (Step-1), the adsorbedethoxy groups on the cations in the TiO₂:Ni matrix are replaced byhydroxyl groups due to the instability of the ethoxide species in thepresence of moisture (Rasko, et al., Applied Catalysis A: General 2004,269, (1-2), 13-25). The hydroxyl clusters of Ti(OH)₄ and Ni(OH)₂ coexistin the aged powder samples, with the Ti(OH)₄ being the more stablecompound due to its low hydroxylation energy (Wang, et al., The Journalof Physical Chemistry A 2010, 114, (28), 7561-7570). These nanopowders,when annealed at 450° C., undergo structural changes owing tothermodynamic considerations (Buonsanti, et al., Chemistry of Materials2013, 25, (8), 1305-1317). In nanocrystals, the self-purificationeffect-high formation energy of defect impurities compared to bulkmaterials-will tend to anneal out dopants during the growth process viadiffusion (Dalpian, et al., Physical review letters 2006, 96, (22),226802; Yang, et al., Journal of the American Chemical Society 2008,130, (46), 15649-15661). At operating temperatures of about 400° C.,specific to the host lattice, the dopant diffusion can take placefollowed by lattice ejection (Chen, et al., Journal of the AmericanChemical Society 2009, 131, (26), 9333-9339). This explains theformation of NiO from the Ni(OH)₂ clusters in the hydrated matrix(Step-2, 3) (Li, et al., Chemical Engineering Journal 2008, 136, (2-3),398-408; Li, et al., Nanoscale 2011, 3, (12), 5103-5109). A similarsequence of processes, i.e., lattice incorporation, lattice diffusion,and lattice ejection has been observed for other doped nanocrystalsystems as well (Xie, et al., Journal of the American Chemical Society2009, 131, (30), 10645-10651). Moreover, the size of the segregated NiOclusters is clearly a function of the concentration of the parentNi(OH)₂ clusters, and therefore, a strong NiO diffraction peak wasobserved for the air exposed powders in the XRD patterns (FIG. 14A).These results demonstrate that the size of the NiO cluster in the hostlattice can be tailored by controlling the moisture interaction with thedopant in the host lattice.

Since the above reaction mechanism indicates that the formation of NiOclusters is an intrinsic process at high doping concentrations, the nextobjective of this work was to control this dopant segregation anddiffusion. Recent literature reports show that the transformation ofamorphous titania into crystalline anatase phase can be achieved within30 min of thermal treatment at temperatures as low as 350° C. (Bhosle,et al., Nanotechnology 2017, 28, (40), 405603; Nikodemski, et al.,Scientific reports 2016, 6, 32830). In order to take advantage of thisrapid heat treatment for locking the dopants in host lattice, TiO₂:Ni(15 mol %) dried powders were rapidly annealed at 450° C. for 30 min.First, FTIR measurements were performed on these NPs to identify anyincombustible species; the spectra showed that lowering of annealingtime from 2 h to 30 min has no effect on the combustion of organicspecies, however, it has an effect on the dopant diffusion. Rapidannealed TiO₂:Ni (15 mol %) NPs exhibited anatase peaks in the XRDpattern, but with poor crystallinity. Moreover, the presence of NiOclusters was not identified in the detailed XRD scans. Therefore, inorder to probe the local crystallinity and the dopant distribution,these NPs were characterized using HRTEM coupled with EELS chemicalmapping (FIG. 39). The HAADF image shown in FIG. 39 indicates mostlyuniform distribution of Ni dopants with some regions of highconcentration. The lattice planes in the diffraction image are alignedalong the (101) plane of anatase TiO₂ (d₁₀₁=3.5 Å) and (001) plane ofNiTiO₃ (d₀₀₁=4.16 Å). From the thermodynamic standpoint, anatase TiO₂ isthe most stable product with the (101) direction being the lowest energyfacet (Yu, et al., Journal of the American Chemical Society 2014, 136,(25), 8839-8842), followed by NiTiO₃ (Kale, Metallurgical and MaterialsTransactions B 1998, 29, (1), 31-38). Furthermore, Ni is eitherdistributed in the lattice of TiO₂ or it is forming a mixed oxide phaseNiTiO₃. These results indicate that the rapid annealing of TiO₂:Ni (15mol %) NPs limits the diffusion of the dopant in the host lattice due tothe short annealing time, and thereby, locking the dopant in the hostlattice or resulting in the formation of a mixed oxide phase. The dopedsystem is therefore quenched in the metastable excited state (NiTiO₃)during the lattice incorporation phase.

Demonstration of Photocatalytic Activity: As a proof-of-concept, thephotocatalytic activity of these TiO₂:Ni (15 mol %) materials wasevaluated using a cationic dye, methylene blue. The difference in thecrystal structure of these NPs demonstrated a difference in thephotodegradation rates.

The phenomenon of limiting the dopant diffusion to form dopant oxideclusters by rapid thermal treatment was not only limited to Ni²⁺, butobserved with other first row TM element, Co²⁺. The XRD pattern ofTiO₂:Co²⁺ (15 mol %) is shown for the rapid and slow annealed NPs (FIG.16A). The slow annealed NPs show the diffraction peaks corresponding toCo₃O₄ and anatase phases, whereas the rapidly annealed NPs crystallizedin doped anatase TiO₂ phase. This unique finding for the highly dopedTiO₂ nanocrystals—that the crystal structure can be tailored by varyingthe annealing rate-will provide an important means to improve theperformance of these host materials in solar—based technologies.

In summary, TiO₂:Ni (15 mol %) NPs were synthesized using sol-gelchemistry. Bulk characterization techniques such as UV-Vis and XRDconfirmed the optical and structural properties of the doped system. Theaging of the dried TiO₂:Ni (15 mol %) NPs prior to annealing wasobserved to impact the segregation of the NiO clusters in TiO₂. Thedifference in the surface hydroxyl concentration for the aged versusnon-aged powders was observed using FTIR measurements. Furthermore,time-resolved UV-Vis absorption measurements on the dried TiO₂:Ni (15mol %) powders showed a systematic increase in the crystal fieldsplitting energy with the aging time, indicating the replacement ofethoxy with hydroxyl groups on surface cations. TGA-DSC studies wereperformed on the aged and non-aged TiO₂:Ni (15 mol %) NPs to elucidatethe moisture-dopant interaction mechanism. Aging the dried powders inair results in the formation of hydroxyl clusters, i.e., Ni(OH)₂, whichcrystallize into NiO clusters upon annealing. The lower calcinationtemperature of NiO and the higher thermodynamic stability ofTi(OH)₄/TiO₂ clusters favored the coexistence of NiO and TiO₂ phases inthe TiO₂:Ni (15 mol %) NPs. Next, rapid annealing was performed on theamorphous TiO₂:Ni (15 mol %) powders to control the dopant segregation.HRTEM and EELS images confirmed the formation of doped anatase andNiTiO₃ phases. This ability to spatially control the dopant localenvironment in a solid host will provide opportunities to tailor theoptoelectronic properties of the material for select applications.

The disclosures of each and every patent, patent application, andpublication cited herein are hereby incorporated herein by reference intheir entirety. While this invention has been disclosed with referenceto specific embodiments, it is apparent that other embodiments andvariations of this invention may be devised by others skilled in the artwithout departing from the true spirit and scope of the invention. Theappended claims are intended to be construed to include all suchembodiments and equivalent variations.

1. A composite material comprising: a rare earth doped core; atransition metal doped shell disposed over the core; and afunctionalizable group on the surface of the transition metal dopedshell.
 2. The composite material of claim 1, wherein the compositematerial forms a shape selected from the group consisting of acore-shell nanoparticle, a nanowire, a nanorod, and a thin film disposedover a substrate.
 3. The composite material of claim 1, wherein thecomposite material is a core-shell nanoparticle.
 4. The compositematerial of claim 1, wherein the rare earth doped core comprisesβ-NaYF₄.
 5. The composite material of claim 1, wherein the rare earthdoped core comprises at least one rare earth selected from the groupconsisting of Er, Yb, Tb, Tm, and Ho.
 6. The composite material of claim1, wherein the transition metal doped shell comprises TiO₂.
 7. Thecomposite material of claim 1, wherein the transition metal doped shellcomprises at least one transition metal selected from the groupconsisting of V, Cr, Mn, Fe, Co, Ni, Cu, and Bi.
 8. The compositematerial of claim 1, wherein the transition metal doped shell comprisesNi.
 9. The composite material of claim 1, wherein the functionalizablegroup is a hydroxide group.
 10. The composite material of claim 1,wherein the transition metal doped shell comprises Bi.
 11. The compositematerial of claim 1, wherein at least one of the rare earth doped coreand the transition metal doped shell comprises YVO₄.
 12. An inkcomprising the composite material of claim
 1. 13. A QR code comprisingthe ink of claim
 12. 14. A method of changing the absorbance or emissionspectrum of a nanoparticle, the method comprising: providing ananoparticle having a rare earth doped core, a transition metal dopedshell, and at least one surface functionalizable group; and treating thenanoparticle with at least one carboxylic acid; wherein the absorbanceor emission spectrum of the nanoparticle is changed upon treatment withthe carboxylic acid.
 15. The method of claim 14, wherein the carboxylicacid is selected from the group consisting ofpara-(fluorosulfonyl)benzoic acid, para-nitrobenzoic acid,para-cyanobenzoic acid, para-bromobenzoic acid, benzoic acid,para-methoxybenzoic acid, and para-aminobenzoic acid.
 16. The method ofclaim 14, wherein the carboxylic acid is para-aminobenzoic acid.
 17. Themethod of claim 14, wherein the step of treating the core-shellnanoparticle with at least one carboxylic acid comprises: treating thenanoparticle with a first carboxylic acid; and treating the nanoparticlewith a second carboxylic acid.
 18. A method of identifying counterfeitcurrency, the method comprising the steps of: providing an ink having acore-shell nanoparticle with surface functionalizable groups; applyingthe ink during the minting of authentic currency; treating a currencysample with a solution comprising a carboxylic acid; and exposing thecurrency sample to UV light; wherein the treatment with carboxylic acidchanges the emission wavelength of the core-shell nanoparticle.
 19. Themethod of claim 18, wherein the core-shell nanoparticle comprises atransition metal doped shell disposed over a rare earth doped core. 20.The method of claim 19, wherein the transition metal doped shellcomprises at least one transition metal selected from the groupconsisting of V, Cr, Mn, Fe, Co, Ni, Cu, and Bi.